Steel plate and method of production of same

ABSTRACT

Low carbon steel plate excellent impact resistance characteristics after carburizing and quenching and after tempering, characterized by having a predetermined chemical composition, an average grain size of carbides of 0.4 μm to 2.0 μm, an area ratio of pearlite of 6% or less, a ratio of a number of carbides at the ferrite grain boundaries to the number of carbides inside the ferrite grains of over 1, and a Vickers hardness of 100HV to 180HV.

TECHNICAL FIELD

The present invention relates to steel plate and a method of productionof the same.

BACKGROUND ART

Steel plate containing, by mass %, carbon in an amount of 0.1 to 0.4% isbeing used as a material for gears, clutches, and other drive systemparts of automobiles by being used press-formed, enlarging holes, bent,drawn, thickened, and thinned and cold forged by combinations of thesame from a blank. Compared with conventional hot forging etc., withcold forging, there is the problem that the amount of strain accumulatedin the material becomes higher, cracks of the material and buckling atthe time of shaping are invited, and deterioration of the partcharacteristics is caused.

In particular, to obtain wear resistance, after the shaped material iscarburized, quenched, and tempered, residual stress is caused by theheat treatment, so formation and growth of fracture from the crackedparts and buckled parts are invited. To use such a part for the drivesystem, an impact resistance characteristic is sought for preventingfracture due to brittleness in the face of the large instantaneous loadapplied to the start of engagement of the gears at the time of startupetc., so excellent cold forgeability and impact resistancecharacteristic after carburizing, quenching, and tempering are beingdemanded from the above steel plate.

Up to now, various proposals have been made regarding arts for improvingthe cold forgeability of steel plate and the impact resistancecharacteristic after carburization (for example, see PLTs 1 to 5).

For example, PLT 1 discloses, as steel for machine structural useimproving toughness by suppressing coarsening of crystal grains incarburization heat treatment, steel for machine structural usecontaining, by mass %, C: 0.10 to 0.30%, Si: 0.05 to 2.0%, Mn: 0.10 to0.50%, P: 0.030% or less, S: 0.030% or less, Cr: 1.80 to 3.00%, Al:0.005 to 0.050%, Nb: 0.02 to 0.10%, and N: 0.0300% or less and having abalance of Fe and unavoidable impurities, having a structure before coldworking comprised of ferrite and pearlite structures, and having anaverage grain size of ferrite grains of 15 μm or more.

PLT 2 discloses, as steel excellent in cold workability and carburizingand quenching ability, steel containing C: 0.15 to 0.40%, Si: 1.00% orless, Mn: 0.40% or less, sol. Al: 0.02% or less, N: 0.006% or less, andB: 0.005 to 0.050%, having a balance of Fe and unavoidable impurities,and having a structure mainly comprised of ferrite phases and graphitephases.

PLT 3 discloses a steel material for carburized bevel gear use excellentin impact strength, a high toughness carburized bevel gear, and a methodof production of the same.

PLT 4 discloses steel for carburized part use having excellentworkability while suppressing coarsening of crystal grains even withsubsequent carburization and having an excellent impact resistancecharacteristic and impact fatigue resistance characteristic in a partproduced by spheroidal annealing, then a cold forging and a carburizing,quenching, and tempering process.

PLT 5 discloses as cold tool steel for plasma carburization use a steelcontaining C: 0.40 to 0.80%, Si: 0.05 to 1.50%, Mn: 0.05 to 1.50%, andV: 1.8 to 6.0%, further containing one or more of Ni: 0.10 to 2.50%, Cr:0.1 to 2.0%, and Mo: 3.0% or less, and having a balance of Fe andunavoidable impurities.

CITATION LIST Patent Literature

PLT 1: Japanese Patent Publication No. 2013-040376A

PLT 2: Japanese Patent Publication No. 06-116679A

PLT 3: Japanese Patent Publication No. 09-201644A

PLT 4: Japanese Patent Publication No. 2006-213951A

PLT 5: Japanese Patent Publication No. 10-158780A

SUMMARY OF INVENTION Technical Problem

The structure of the steel for machine structural use of PLT 1 is astructure of ferrite+pearlite. This structure, compared with aferrite+cementite structure, has a large hardness, so wear of the die incold forging cannot be suppressed and the steel cannot necessarily besaid to be steel for machine structural use excellent in coldforgeability.

In the steel of PLT 2, the graphitization treatment of the cementiterequires annealing at a high temperature. A drop in the yield and anincrease in the manufacturing costs cannot be suppressed.

The method of production of PLT 3 requires further hot forging aftercold forging and carburizing. Since hot forging is essential, this isnot a method of production leading to fundamentally lower costs.

It is unclear if the steel for carburized part use of PLT 4 can exhibitssimilar effects in cold forging given a large strain. Furthermore, thespecific form of the structure and method of control of the structureare also unclear, so this cannot be said to be steel exhibitingexcellent workability even in the plate forging growing in use in recentyears and other shaping by forging cold while giving a large strain.

PLT 5 does not disclose at all the findings and art relating to theoptimum components and form of structure for improving the formabilityof steel, in particular cold forgeability.

The present invention, in consideration of the above prior art, has asits problem the provision of steel plate excellent in cold forgeabilityand impact resistance characteristic after carburizing, quenching, andtempering, in particular suitable for obtaining a high cycle gear orother part by forming a plate and of a method of production of the same.

Solution to Problem

To solve the above problem and obtain steel plate suitable for amaterial such as a drive system part, it is understood that in a steelplate containing the C required for raising the hardenability,enlargement of the ferrite in grain size, spheroidization of thecarbides (mainly cementite) to a suitable grain size, and reduction ofthe pearlite structures are preferable. This is due to the followingreasons.

A ferrite phase is low in hardness and high in ductility. Therefore, ina structure mainly comprised of ferrite, it becomes possible to increasethe grain size so as to raise the formability of the material.

Carbides, by being made to suitably disperse in the metal structure, canmaintain the formability of the material while imparting an excellentwear resistance and rolling fatigue characteristic, so provides astructure essential for drive system parts. Further, the carbides in thesteel plate are strong particles obstructing slip.

By forming carbides at the ferrite grain boundaries, it is possible toprevent propagation of slip exceeding the crystal grain boundaries andsuppress the formation of shear zones. Thus the cold forgeability isimproved and, simultaneously, the formability of steel plate is alsoimproved.

However, cementite is a hard, brittle structure. If a laminar structurewith ferrite present, that is, in the state of pearlite, the steelbecomes hard and brittle, so it has to be present in a spheroidal form.If considering the cold forgeability and the occurrence of fractures atthe time of forging, its grain size has to be a suitable range.

However, no method of production for realizing the above structure hasbeen disclosed up to now. Therefore, the inventors intensivelyresearched a method of production for realizing the above structure.

As a result, they discovered the following: To make the metal structureof the steel plate after coiling after hot rolling a bainite structureof fine pearlite or fine ferrite with small lamellar spacing in whichcementite is dispersed, the steel plate is coiled at a relatively lowtemperature (400° C. to 550° C.). By coiling at a relatively lowtemperature, the cementite dispersed in the ferrite also easily becomesspheroidal. Next, the cementite is partially made spheroidal byannealing at a temperature just under the Ac1 point as first stageannealing. Next, as second stage annealing, part of the ferrite grainsis left while part is transformed to austenite by annealing at atemperature between the Ac1 point and Ac3 point (so-called dual phaseregion of ferrite and austenite). By then making the remaining ferritegrains grow while slowly cooling the steel while using these as nucleito transform the austenite to ferrite, it is possible to obtain largeferrite phases and make cementite precipitate at the grain boundaries torealize the above structure.

That is, the method of production of steel plate simultaneouslysatisfying hardenability and formability is difficult to realize even ifdesigning the hot rolling conditions, annealing conditions, etc. assingle processes. It was discovered that this can be realized byoptimization by a so-called integral process of hot rolling, anannealing process, etc.

Further, improvement of the drawability at the time of cold forgingrequires the reduction of plastic anisotropy. It was discovered that forsuch improvement, adjustment of the hot rolling conditions is important.

The present invention was made based on these discoveries and has as itsgist the following:

(1) A steel plate being low carbon steel plate having a chemicalcomposition containing, by mass %, C: 0.10 to 0.40%, Si: 0.01 to 0.30%,Mn: 0.30 to 1.00%, Al: 0.001 to 0.10%, Cr: 0.50 to 2.00%, Mo: 0.001 to1.00%, P: 0.020% or less, S: 0.010% or less, N: 0.020% or less, O:0.020% or less, Ti: 0.010% or less, B: 0.0005% or less, Sn: 0.050% orless, Sb: 0.050% or less, As: 0.050% or less, Nb: 0.10% or less, V:0.10% or less, Cu: 0.10% or less, W: 0.10% or less, Ta: 0.10% or less,Ni: 0.10% or less, Mg: 0.050% or less, Ca: 0.050% or less, Y: 0.050% orless, Zr: 0.050% or less, La: 0.050% or less, and Ce: 0.050% and havinga balance of Fe and impurities, the metal structure of the steel platehaving a carbide grain size of 0.4 to 2.0 μm, a pearlite area ratio of6% or less, and a ratio of a number of carbides at the ferrite grainboundaries to the number of carbides inside the ferrite grains of over1, the steel plate having a Vickers hardness of 100 HV to 180 HV.

(2) A method of production of the steel plate according to (1), themethod of production comprising the steps of: hot rolling a steel slabof a chemical composition according to claim 1, completing finish hotrolling in a 650° C. to 950° C. temperature region to obtain a hotrolled steel plate; coiling the hot rolled steel plate at 400° C. to600° C.; pickling the coiled hot rolled steel plate and heating thepickled hot rolled steel plate by a 30° C./hour to 150° C./hour heatingrate to a 650° C. to 720° C. annealing temperature and holding it therefor 3 hours to 60 hours as first stage annealing; then heating the hotrolled steel plate to an annealing temperature of 725° C. to 790° C. bya heating rate of 1° C./hour to 80° C./hour and holding the steel platefor 3 hours to 50 hours as second stage annealing; and cooling theannealed hot rolled steel plate to 650° C. by a cooling rate of 1°C./hour to 100° C./hour.

According to the present invention, it is possible to provide a steelplate excellent in cold forgeability and impact resistancecharacteristic after carburizing, quenching, and tempering, inparticular one suitable for obtaining a high cycle gear or other part byforming a plate.

BRIEF DESCRIPTION OF DRAWINGS

FIGS. 1A to 1C are views schematically showing a summary of the coldforging test and the form of a crack introduced by cold forging. FIG. 1Ashows a disk-shaped test material cut out from a hot rolled steel plate,while FIG. 1B shows the shape of a test material after cold forging,while FIG. 1C shows the cross-sectional shape of the test material aftercold forging.

FIG. 2 is a view schematically showing a summary of a drop weight testevaluating the impact resistance characteristic of a sample performingcarburizing, quenching, and tempering.

FIG. 3 is a view showing a relationship among a ratio of a number ofcarbides at the grain boundaries to the number of carbides in thegrains, the crack length of the cold forging test piece, and the impactresistance characteristic after carburizing, quenching, and tempering.

FIG. 4 is a view showing another relationship among a ratio of a numberof carbides at the grain boundaries to the number of carbides in thegrains, the crack length of the cold forging test piece, and the impactresistance characteristic after carburizing, quenching, and tempering.

DESCRIPTION OF EMBODIMENTS

Below, the present invention will be explained in detail. First, thereasons for limitation of the chemical composition of the steel plate ofthe present invention will be explained. Here, the “%” according to thechemical composition means “mass %”.

C: 0.10 to 0.40%

C is an element forming carbides in steel and effective forstrengthening the steel and refining the ferrite grains. To suppress theformation of a matte surface in cold working and secure surface beautyof a cold forged part, suppression of coarsening of the ferrite grainsize is essential, but if less than 0.10%, the carbides becomeinsufficient in volume fraction and coarsening of the carbides duringannealing can no longer be suppressed, so C is made 0.10% or more.Preferably it is 0.11% or more.

On the other hand, if exceeding 0.40%, the carbides increase in volumefraction, a large amount of cracks are formed acting as starting pointsof breakage at the time of an instantaneous load and a drop in theimpact resistance characteristic is invited, so C is made 0.40% or less.Preferably it is 0.38% or less.

Si: 0.01 to 0.30%

Si is an element which acts as a deoxidizing agent and further has aneffect on the form of the carbides. To reduce the number of carbides inthe ferrite grains giving the deoxidizing effect and increase the numberof carbides at the ferrite grain boundaries, it is necessary to usetwo-stage step type annealing to produce austenite phases duringannealing, make the carbides dissolve once, then gradually cool thestructure to promote the formation of carbides at the ferrite grainboundaries.

If Si exceeds 0.30%, the ferrite falls in ductility, fractures areeasily formed at the time of cold forging, and the cold forgeability andimpact resistance characteristic after carburizing, quenching, andtempering deteriorate, so Si is made 0.30% or less. Preferably it is0.28% or less.

Si is preferably as low as possible, but reduction to less than 0.01%invites a large increase in refining costs, so Si is made 0.01% or more.Preferably it is 0.02% or more.

Mn: 0.30 to 1.00%

Mn is an element controlling the form of carbides in two-stage step typeannealing. If less than 0.30%, in the gradual cooling after second stageannealing, it becomes difficult to form carbides at the ferrite grainboundaries, so Mn is made 0.30% or more. Preferably it is 0.33% or more.

On the other hand, if exceeding 1.00%, the toughness after carburizing,quenching, and tempering falls, so Mn is made 1.00% or less. Preferablyit is 0.96% or less.

Al: 0.001 to 0.10%

Al is an element acting as a deoxidizing agent of steel and stabilizingferrite. If less than 0.001%, the effect of addition is not sufficientlyobtained, so Al is made 0.001% or more. Preferably it is 0.004% or more.

On the other hand, if exceeding 0.10%, the number ratio of carbides atthe grain boundaries is lowered and an increase in crack length at thetime of cold forging is invited, so Al is made 0.10% or less. Preferablyit is 0.09% or less.

Cr: 0.50 to 2.00%

Cr and Mo are elements which improve the toughness. Cr is an elementeffective for stabilization of carbides at the time of heat treatment.If less than 0.50%, it becomes difficult to cause carbides to remain atthe time of carburization, coarsening of the austenite grain size at thesurface layer is invited, and a drop in the impact resistancecharacteristic is caused, so Cr is made 0.50% or more. Preferably it is0.52% or more.

On the other hand, if exceeding 2.00%, the amount of Cr concentrating atthe carbides increases and a large amount of fine carbides remain in theaustenite phases produced by the two-stage step type annealing, carbidesremain in the grains after gradual cooling, the hardness increases andnumber ratio of carbides at the grain boundaries fall and the coldforgeability falls, so Cr is made 2.00% or less. Preferably it is 1.94%or less.

Mo: 0.001 to 1.00%

Mo is an element effective for control of the form of carbides. If lessthan 0.001%, the effect of addition is not sufficiently obtained, so Mois made 0.001% or more. Preferably it is 0.017% or more.

On the other hand, if exceeding 1.00%, Mo concentrates in the carbidesand stable carbides increase in the austenite phase as well, so aftergradual cooling, carbides are present in the grains as well, an increasein hardness and drop in number ratio of carbides at the grain boundariesare invited, and the cold forgeability falls, so Mo is made 1.00% orless. Preferably it is 0.94% or less.

The following elements are impurities and have to be controlled tocertain amounts or less.

P: 0.020% or Less

P is an element segregating at the ferrite grain boundaries andsuppressing the formation of carbides at the grain boundaries. Thesmaller amount is preferable. The content of P may also be 0, but a longtime is required for refining in order to make the purity a high one ofless than 0.0001% in a refining process and a large increase in themanufacturing cost is invited, so the de facto lower limit is 0.0001 to0.0013%.

On the other hand, if exceeding 0.020%, the number ratio of carbides atthe grain boundaries falls and the cold forgeability falls, so P is made0.020% or less. Preferably it is 0.018% or less.

S: 0.010% or Less

S is an impurity element forming MnS and other nonmetallic inclusions.The nonmetallic inclusions form starting points of formation offractures at the time of cold forging, so the smaller the S, the better.The content of S may also be 0, but to lower S to less than 0.0001%, therefining costs greatly increase, so the de facto lower limit is 0.0001to 0.0012%.

On the other hand, if exceeding 0.010%, an increase is invited in thecrack length at the time of cold forging, so S is made 0.010% or less.Preferably it is 0.009% or less.

N: 0.020% or Less

N is an element segregating at the ferrite grain boundaries andsuppressing the formation of carbides at the grain boundaries. Thesmaller amount is preferable. The content of N may also be 0, but ifreducing it to less than 0.0001%, the refining costs greatly increase,so the de facto lower limit is 0.0001 to 0.0006%.

On the other hand, if exceeding 0.020%, even if performing dual phaseregion annealing and gradual cooling, the ratio of the number ofcarbides at the ferrite grain boundaries with respect to the number ofcarbides in the ferrite grains becomes less than 1 and the coldforgeability fall, so N is made 0.020% or less. Preferably it is 0.017%or less.

O: 0.0001 to 0.020%

O is an element forming oxides in the steel. The oxides present in theferrite grains become sites for production of carbides, so the smallerthe amount, the better. The content of 0 may also be 0, but if reducing0 to less than 0.0001%, the refining costs greatly increase, so the defacto lower limit is 0.0001 to 0.0006%.

On the other hand, if exceeding 0.020%, the ratio of the number ofcarbides at the ferrite grain boundaries with respect to the number ofcarbides in the ferrite grains becomes less than 1 and the coldforgeability falls, so 0 is made 0.020% or less. Preferably it is 0.017%or less.

Ti: 0.010% or Less

Ti is an element important for control of the form of the carbides. Itis an element by which, by inclusion in a large amount, formation ofcarbides in the ferrite grains is promoted. The smaller amount ispreferable. The content of Ti may also be 0, but if reducing it to lessthan 0.0001%, the refining costs greatly increase, so the de facto lowerlimit is 0.0001 to 0.0006%.

On the other hand, if over 0.010%, the ratio of the number of carbidesat the ferrite grain boundaries to the number of carbides inside theferrite grains becomes less than 1 and the cold forgeability falls, soTi is made 0.010% or less. Preferably it is 0.007% or less.

B: 0.0005% or Less

B is an element effective for control of slip of dislocations at thetime of cold forging. By inclusion of a large amount, activity of theslip system is limited, so the smaller the amount of B, the better. Thecontent of B may also be 0. Fine care is required for detection of lessthan 0.0001% of B. Depending on the analysis device, it is below thelower limit of detection.

On the other hand, if exceeding 0.0005%, cross slip of dislocations atthe shear zone formed by the cold forging is suppressed. Strainconcentrates locally and fractures are formed, so B is made 0.0005% orless. Preferably it is 0.0005% or less.

Sn: 0.050% or Less

Sn is an element entering from the steel starting materials (scraps).The smaller amount is preferable. The content of Sn may also be 0, butif reducing it to less than 0.001%, the refining costs greatly increase,so the de facto lower limit is 0.001 to 0.002%.

On the other hand, if exceeding 0.050%, the ferrite becomes brittle andthe cold forgeability falls, so Sn is made 0.050% or less. Preferably,it is 0.048% or less.

Sb: 0.050% or Less

Sb, like Sn, is an element entering from the steel starting materials(scraps). Sb segregates at the grain boundaries and lowers the numberratio of carbides at the grain boundaries, so the smaller the amount,the better. The content of Sb may also be 0, but if reducing it to lessthan 0.001%, the refining costs greatly increase, so the de facto lowerlimit is 0.001 to 0.002%.

On the other hand, if exceeding 0.050%, the cold forgeability falls, soSb is made 0.050% or less. Preferably, it is 0.048% or less.

As: 0.050% or Less

As is an element which enters from the steel starting materials (scraps)like Sn and Sb. As segregates at the grain boundaries and lowers thenumber ratio of carbides at the grain boundaries, so the content ispreferably small. The content of As may also be 0, but if reducing it toless than 0.001%, the refining cost greatly increases, so the de factolower limit is 0.001 to 0.002%.

On the other hand, if over 0.050%, the number ratio of carbides at thegrain boundaries falls and the cold forgeability falls, so As is made0.050% or less. Preferably it is 0.045% or less.

The steel plate of the present invention has the above elements as basicelements, but may further contain the following elements for the purposeof improving the cold forgeability and other characteristics. Thefollowing elements are not essential for obtaining the effects of thepresent invention, so the contents may also be 0.

Nb: 0.10% or Less

Nb is an element effective for control of the form of the carbides.Further, it is an element refining the structure and contributing toimprovement of the toughness. If less than 0.001%, the effect ofaddition is not sufficiently obtained, so Nb is preferably made 0.001%or more. More preferably, it is 0.002% or more.

On the other hand, if over 0.10%, a large number of fine Nb carbidesprecipitate, the strength excessively rises, and, further, the numberratio of carbides at the grain boundaries falls and the coldforgeability falls, so Nb is made 0.10% or less. Preferably it is 0.09%or less.

V: 0.10% or Less

V, like Nb, is an element effective for control of the form of thecarbides. Further, it is an element refining the structure andcontributing to improvement of the toughness. If less than 0.001%, theeffect of addition is not sufficiently obtained, so V is preferably made0.001% or more. More preferably, it is 0.004% or more.

On the other hand, if over 0.10%, a large number of fine V carbidesprecipitate, the strength excessively rises, and, further, the numberratio of carbides at the grain boundaries falls and the coldforgeability falls, so V is made 0.10% or less. Preferably, it is 0.09%or less.

Cu: 0.10% or Less

Cu is an element forming fine precipitates and contributing toimprovement of the strength. If less than 0.001%, the effect ofimprovement of the strength is not sufficiently obtained, so Cu ispreferably made 0.001% or more. More preferably, it is 0.008% or more.

On the other hand, if over 0.10%, red hot embrittlement occurs duringhot rolling and the productivity falls, so Cu is made 0.10% or less.Preferably, it is 0.09% or less.

W: 0.10% or Less

W, like Nb and V, is an element effective for control of the form of thecarbides. If less than 0.001%, the effect of addition is notsufficiently obtained, so W is preferably made 0.001% or more. Morepreferably, it is 0.003% or more.

On the other hand, if over 0.10%, a large number of fine W carbidesprecipitate, the strength excessively rises, and, further, the numberratio of carbides at the grain boundaries falls and the coldforgeability falls, so W is made 0.10% or less. Preferably, it is 0.08%or less.

Ta: 0.10% or Less

Ta, like Nb, V, and W, is an element effective for control of the formof the carbides. If less than 0.001%, the effect of addition is notsufficiently obtained, so Ta is preferably made 0.001% or more. Morepreferably, it is 0.007% or more.

On the other hand, if over 0.10%, a large number of fine Ta carbidesprecipitate, the strength excessively rises, and, further, the numberratio of carbides at the grain boundaries falls and the coldforgeability falls, so Ta is made 0.10% or less. Preferably, it is 0.09%or less.

Ni: 0.10% or Less

Ni is an element effective for improvement of the impact resistancecharacteristic of parts. If less than 0.001%, the effect of addition isnot sufficiently obtained, so Ni preferably is made 0.001% or more. Morepreferably it is 0.002% or more.

On the other hand, if over 0.10%, the number ratio of carbides at thegrain boundaries falls and the cold forgeability falls, so Ni is made0.10% or less. Preferably, it is 0.09% or less.

Mg: 0.050% or Less

Mg is an element which can control the form of sulfides by addition in atrace amount. If less than 0.0001%, the effect of addition is notsufficiently obtained, so Mg preferably is made 0.0001% or more. Morepreferably it is 0.0008% or more.

On the other hand, if over 0.050%, the ferrite becomes brittle and thecold forgeability falls, so Mg is made 0.050% or less. Preferably it is0.049% or less.

Ca: 0.050% or Less

Ca, like Mg, is an element which can control the form of sulfides byaddition in a trace amount. If less than 0.001%, the effect of additionis not sufficiently obtained, so Ca preferably is made 0.001% or more.More preferably it is 0.003% or more.

On the other hand, if over 0.050%, coarse Ca oxides are formed andbecome starting points of fracture at the time of cold forging, so Ca ismade 0.050% or less. Preferably it is 0.04% or less.

Y: 0.050% or Less

Y, like Mg and Ca, is an element which can control the form of sulfidesby addition in a trace amount. If less than 0.001%, the effect ofaddition is not sufficiently obtained, so Y preferably is made 0.001% ormore. More preferably it is 0.003% or more.

On the other hand, if over 0.050%, coarse Y oxides are formed and becomestarting points of fracture at the time of cold forging, so Y is made0.050% or less. Preferably it is 0.031% or less.

Zr: 0.050% or Less

Zr, like Mg, Ca, and Y, is an element which can control the form ofsulfides by addition in a trace amount. If less than 0.001%, the effectof addition is not sufficiently obtained, so Zr preferably is made0.001% or more. More preferably it is 0.004% or more.

On the other hand, if over 0.050%, coarse Zr oxides are formed andbecome starting points of fracture at the time of cold forging, so Zr ismade 0.050% or less. Preferably it is 0.045% or less.

La: 0.050% or Less

La is an element effective for control of the form of sulfides byaddition in a trace amount. Further, it is an element which segregatesat the grain boundaries and lowers the number ratio of carbides at thegrain boundaries. If less than 0.001%, the effect of control of the formis not sufficiently obtained, so La is preferably made 0.001% or more.More preferably, it is 0.003% or more.

On the other hand, if over 0.050%, the number ratio of carbides at thegrain boundaries falls and the cold forgeability falls, so La is made0.050% or less. Preferably it is 0.047% or less.

Ce: 0.050% or Less

Ce, like La, is an element able to control the form of sulfides byaddition in a trace amount. Further, it is an element which segregatesat the grain boundaries and lowers the number ratio of carbides at thegrain boundaries. If less than 0.001%, the effect of control of the formis not sufficiently obtained, so Ce is preferably made 0.001% or more.More preferably, it is 0.003% or more.

On the other hand, if exceeding 0.050%, the number ratio of carbides atthe grain boundaries falls and the cold forgeability falls, so Ce ismade 0.050% or less. Preferably it is 0.046% or less.

Note that, the remainder of the chemical composition of the steel plateof the present invention is comprised of Fe and unavoidable impurities.

Next, the structure of the steel plate of the present invention will beexplained.

The structure of the steel plate of the present invention issubstantially a structure comprised of ferrites and carbides. Thecarbides include cementite (Fe₃C) which is a compound of iron andcarbon, a compound obtained by substituting Mn, Cr, etc. for the Featoms in the cementite, and alloy carbides (M₂₃C₆, M₆C, MC, etc., whereM is Fe and other metal elements).

When forming steel plate into a predetermined part shape, a shear zoneis formed at the macrostructure of the steel plate and slip deformationoccurs concentrated near the shear zone. In slip deformation, along withproliferation of dislocations, a region of a high dislocation density isformed near the shear zone. Along with the increase in the amount ofstrain imparted to the steel plate, slip deformation is promoted and thedislocation density increases.

In cold forging, strong working is performed with an equivalent strainexceeding 1. For this reason, in conventional steel plate, it was notpossible to prevent the formation of voids and/or cracks along with theincrease in dislocation density and was difficult to improve the coldforgeability.

To solve this difficult problem, it is effective to suppress theformation of a shear zone at the time of forming. From the viewpoint ofthe microstructure, formation of a shear zone can be understood as thephenomenon of slip occurring at a certain one grain crossing the crystalgrain boundary and being continuously propagated to the adjoining grain.Accordingly, to suppress the formation of a shear zone, it is necessaryto prevent propagation of slip crossing crystal grain boundaries.

The carbides in steel plate are strong particles inhibiting slip. Byforming carbides at the ferrite grain boundaries, it becomes possible tosuppress the formation of a shear zone and improve the coldforgeability.

To obtain such an effect, carbides have to be made to disperse in themetal structure in suitable sizes. Therefore, the average particle sizeof carbides is made 0.4 μm to 2.0 μm. If the particle size of thecarbides is less than 0.4 μm, the steel plate remarkably increases inhardness and the cold forgeability falls. More preferably it is 0.6 μmor more.

On the other hand, if the average particle size of the carbides exceeds2.0 μm, at the time of cold forming, the carbides form starting pointsof fractures. More preferably, it is 1.95 μm or less.

Further, cementite, a carbide of iron, has a hard and brittle structure.If present in the form of pearlite, which is a layered structure withferrite, the steel becomes hard and brittle. Therefore, pearlite has tobe reduced as much as possible. In the steel plate of the presentinvention, the area ratio is made 6% or less.

Pearlite has a unique lamellar structure, so can be discerned byobservation by an SEM or optical microscope. By calculating the regionof the lamellar structure at any cross-section, the area ratio of thepearlite can be found.

Based on theory and principle, cold forgeability is considered to bestrongly affected by the rate of coverage of the ferrite grainboundaries by carbides. High precision measurement is sought, butmeasurement of the rate of coverage of ferrite grain boundaries bycarbides in a three-dimensional space requires serial sectioning SEMobservation using an FIB to repeatedly cut a sample for observation in ascanning electron microscope or 3D EBSP observation. A massivemeasurement time is required and technical knowhow has to be built up.

The inventors clarified this and searched for a simpler, higherprecision indicator for evaluation and as a result discovered that it ispossible to evaluate the cold forgeability by using the ratio of thenumber of carbides at the ferrite grain boundaries to the number ofcarbides in the ferrite grains as an indicator and that if the ratio ofthe number of carbides at the ferrite grain boundaries to the number ofcarbides in the ferrite grains exceeds 1, the cold forgeabilityremarkably rises.

Note that, buckling, folding, and twisting of the steel plate occurringat the time of cold forging occur due to localization of strainaccompanying the formation of a shear zone, so similarly by formingcarbides at the ferrite grain boundaries to reduce the formation of ashear zone and localization of strain, it is possible to suppressbuckling, folding, and twisting.

The carbides are observed by a scanning electron microscope. Beforeobservation, the sample for observation of the structure is polished bywet polishing by Emery paper and a diamond abrasive having an averageparticle size of 1 μm, the observed surface is polished to a mirrorfinish, then a saturated picric acid-alcohol solution is used to etchthe structure.

The magnification of the observation was made 3000× and images of eightfields of 30 μm×40 μm at a plate thickness ¼ layer were captured atrandom. The obtained structural images were analyzed by image analyzingsoftware such as one made by Mitani Shoji (Win ROOF) to measure indetail the areas of the carbides contained in those regions. The circleequivalent diameters (=2×·(area/3.14)) were found from the areas of thecarbides and the average value was made the particle size of thecarbides.

Note that, to keep down the effect of measurement error due to noise,carbides with an area of 0.01 μm² or less are excluded from the coverageof the evaluation.

The number of carbides which present at the ferrite grain boundaries arecounted, the number of carbides at the grain boundaries are subtractedfrom the total number of carbides, and the number of carbides in theferrite grains are found. Based on the measured number, the number ratioof carbides at the grain boundaries with respect to the carbides insidethe ferrite grains is calculated.

By making as the structure after annealing a structure with ferritegrains of a size of 3.0 μm to 50.0 μm, it is possible to improve thecold forgeability. If the size of the ferrite grains is less than 3 μm,the hardness increases and fractures and cracks easily form at the timeof cold forging, so the ferrite grain size is preferably 3.0 μm or more.More preferably it is 7.5 μm or more.

On the other hand, if the ferrite grain size is over 50.0 μm, the numberof carbides on the crystal grain boundaries suppressing the propagationof slip is decreased and the cold forgeability falls, so the ferritegrain size is preferably 50.0 μm or less. More preferably it is 37.9 μmor less.

The ferrite grain size is measured by using the above-mentionedprocedure to polish the observed surface of the sample for observationof structure to a mirror finish, then observing the structure of theobserved surface etched by a 3% nitric acid-alcohol solution by anoptical microscope or scanning electron microscope and applying the linesegment method to the captured image.

By making the Vickers hardness of the steel plate 100 HV to 180 HV, itis possible to improve the cold forgeability and impact resistancecharacteristic after carburizing, quenching, and tempering. If theVickers hardness is less than 100 HV, buckling easily occurs during coldforging, folding and twisting occur at the buckled part, and the impactresistance characteristic falls, so the Vickers hardness is made 100 HVor more. Preferably it is 110 HV or more.

On the other hand, if the Vickers hardness exceeds 180 HV, the ductilityfalls, internal cracking easily occurs at the time of cold forging, andthe impact resistance characteristic deteriorates, so the Vickershardness is made 180 HV or less. Preferably it is 170 HV or less.

Next, the method of evaluation of the cold forgeability will beexplained.

FIGS. 1A to 1C schematically show an outline of the cold forging testand form of a crack introduced by cold forging. FIG. 1A shows adisk-shaped test material cut out from a hot rolled steel plate, FIG. 1Bshows the shape of a test material after cold forging, and FIG. 1C showsthe cross-sectional shape of a cold forged test material.

As shown in FIGS. 1A to 1C, from a plate thickness 5.2 mm hot rolledsteel plate, a diameter 70 mm disk-shaped test material 1 was cut out(see FIG. 1A) and deep drawn so as to prepare a cup-shaped test materialwith a bottom surface of a diameter of 30 mm (not shown). Next, a oneshot forming press made by Mori Ironworks was used to thicken thevertical wall parts of a cup-shaped test material by a thickening ratioof 1.54 (=8 mm/5.2 mm) (cold forging) to prepare a cup-shaped testmaterial 2 with a diameter of 30 mm, a height of 30 mm, and a verticalwall thickness of 8 mm (see FIG. 1B).

The thickened cup-shaped test material 2 is cut by a wire cut electricaldischarge machine made by FANUC so that the cross-section of thediameter part appeared (see FIG. 1C. The cut surface is polished to amirror finish and the presence of a fracture 3 at the cut surface wasconfirmed. The ratio of the maximum length of fracture L present at thevertical wall parts to the thickness of the vertical wall part afterthickening (=maximum length of fracture L/thickness of vertical wallpart after thickening 8 mm) is measured. This measured value is used toevaluate the cold forgeability.

Note that, even if the initial plate thickness is other than 5.2 mm, ifadjusting the diameter of the cut out disk-shaped test material so thatthe height of the vertical wall after thickening becomes 30 mm andforming the material by a thickening ratio of the same 1.54, it ispossible to reproduce the results of evaluation without regard as to theinitial plate thickness, so the hot rolled steel plate covered by thepresent invention is not limited to a plate thickness 5.2 mm hot rolledsteel plate. The present invention can improve the cold forgeability andthe impact resistance characteristic after carburizing, quenching, andtempering even in a general plate thickness (2 to 15 mm) hot rolledsteel plate.

Next, the method of production of the present invention will beexplained. The technical idea of the method of production of the presentinvention is to integrally manage the rolling conditions and annealingconditions when producing steel plate from a steel slab of theabove-mentioned chemical composition so as to improve the coldforgeability and the impact resistance characteristic after carburizing,quenching, and tempering.

The features of the method of production of the present invention willbe explained next.

Features of Hot Rolling

Molten steel having the required chemical composition is continuouslycast into a slab. The slab is used for hot rolling as is in accordancewith an ordinary method or is cooled once, then heated and used for hotrolling. The finish hot rolling is ended in the 650° C. to 950° C.temperature region. The hot rolled steel plate after finish rolling iscooled on the ROT and coiled by a coiling temperature of 400° C. to 600°C.

Features of Annealing

The hot rolled steel plate is pickled, then held at two temperatureregions as two-stage step type annealing, but at that time, in the firststage annealing, the hot rolled steel plate is heated until theannealing temperature by a 30° C./hour to 150° C./hour heating rate andheld at a 650° C. to 720° C. temperature region for 3 hours to 60 hoursfor annealing.

At the next second stage annealing, the hot rolled steel plate is heateduntil the annealing temperature by a 1° C./hour to 80° C./hour heatingrate and held at a 725° C. to 790° C. temperature region for 3 hours to50 hours for annealing.

Next, the annealed hot rolled steel plate is cooled down to 650° C. by acooling rate of 1° C./hour to 100° C./hour, then is cooled down to roomtemperature.

Due to the linkage between these hot rolling conditions and annealingconditions, low carbon steel plate excellent in cold forgeability andimpact resistance characteristic after carburizing, quenching, andtempering can be obtained.

Below, the conditions of steps of the method of production of thepresent invention will be specifically explained.

Hot Rolling

Finish hot rolling temperature: 650° C. to 950° C.

Coiling temperature: 400° C. to 600° C.

Molten steel having the required chemical composition is continuouslycast into a slab. The slab is used for hot rolling as is or cooled once,then heated. The finish hot rolling is ended in the 650° C. to 950° C.temperature region. The hot rolled steel plate is coiled at 400° C. to600° C.

The slab heating temperature is preferably 1300° C. or less, while theheating time where the slab is held at a temperature of the slab surfacelayer of 1000° C. or more is preferably 7 hours or less.

If the heating temperature exceeds 1300° C. or the heating time exceeds7 hours, the decarburization of the slab surface layer becomesremarkable. At the time of heating before hardening, the austenitegrains of the surface layer abnormally grow and the impact resistancecharacteristic falls, so the heating temperature is preferably 1300° C.or less and the heating time is preferably 7 hours or less. Morepreferably, the heating temperature is 1280° C. or less, while theheating time is 6 hours or less.

The finish hot rolling is ended at 650° C. to 950° C. in temperature. Ifthe finish hot rolling temperature is less than 650° C., the rollingload remarkably rises due to the increase of the deformation resistanceof the steel material. Furthermore, the amount of roll wear increasesand the productivity falls, so the finish hot rolling temperature ismade 650° C. or more. Preferably it is 680° C. or more.

On the other hand, if the finish hot rolling temperature exceeds 950°C., bulky scale is formed during passage through the ROT (Run OutTable). Due to the scale, flaws form at the surface of the steel plate.At the time of cold forging and/or at the time when an impact load isapplied after carburizing, quenching, and tempering, cracks formstarting from the flaws and the impact resistance characteristic falls,so the finish hot rolling temperature is made 950° C. or less.Preferably it is 920° C. or less.

The cooling rate when cooling the hot rolled steel plate on the ROT ispreferably 10° C./sec to 100° C./sec. If the cooling rate is less than10° C./sec, in the middle of the cooling, it is not possible to suppressthe formation of bulky scale and the formation of flaws due to the sameand the impact resistance characteristic falls, so the cooling rate ispreferably 10° C./sec or more. More preferably it is 20° C./sec or more.

On the other hand, if cooling the hot rolled steel plate from thesurface layer to the inside part of the steel plate by a cooling rate ofover 100° C./sec, the outermost layer part is excessively cooled andbainite, martensite, and other low temperature transformed structuresare formed at the outermost layer part.

After coiling, when the 100° C. to room temperature hot rolled steelplate is paid out, fine cracks form at low temperature transformedstructures. It is difficult to remove the cracks in the followingpickling process and cold rolling process. At the time of cold forgingand/or at the time when an impact load is applied after carburizing,quenching, and tempering, cracks grow starting from those cracks andinvite a drop in the impact resistance characteristic, so the coolingrate is preferably 100° C./sec or less. More preferably it is 80° C./secor less.

Note that, the cooling rate indicates the cooling ability received fromthe cooling facilities in a water spray section at the time when beingcooled on the ROT down to the target temperature of coiling from thetime when the hot rolled steel plate after finish hot rolling is watercooled at a water spray section after passing through a non-water spraysection. It does not show the average cooling rate from the startingpoint of water spray to the temperature at which the steel plate iscoiled up by the coiler.

The coiling temperature is made 400° C. to 600° C. This is a temperaturelower than the general coiling temperature. By coiling the hot rolledsteel plate produced in the above-mentioned condition in thistemperature range, the structure of the steel plate can be made abainite structure in which carbides are dispersed in fine ferrite.

If the coiling temperature is less than 400° C., the austenite, whichwas not transformed before coiling, transforms to hard martensite. Atthe time of discharging the coiled hot rolled steel plate, cracks format the surface layer and the impact resistance characteristics fall, sothe coiling temperature is made 400° C. or more. Preferably, it is 430°C. or more.

On the other hand, if the coiling temperature exceeds 600° C., pearlitewith a large lamellar spacing is formed, high thermal stability bulkyneedle shaped carbides are formed, and needle shaped carbides remaineven after two-stage step type annealing. At the time of cold forging,cracks occur and grow starting from these needle shaped carbides, so thecoiling temperature is made 600° C. or less. Preferably it is 570° C. orless.

The hot rolled steel plate produced under the above conditions waspickled, then held in two temperature regions for two-stage step typeannealing. By treating the hot rolled steel plate by two-stage step typeannealing, the carbides are controlled in stability and the formation ofcarbides at the ferrite grain boundaries is promoted.

First, the technical idea of two-stage step type annealing will beexplained.

By performing the first stage annealing in a temperature region of theAc1 point or less, the carbides are made to coarsen and added metalelements are made to concentrate to raise the thermal stability of thecarbides. After that, the steel is raised to the Ac1 or more intemperature to form austenite in the structure, the fine carbides in theferrite grains are made to dissolve into the austenite, and coarsecarbides are left in the austenite.

By the subsequent gradual cooling, the austenite is transformed toferrite and raises the concentration of carbon in the austenite. Alongwith gradual cooling, carbon atoms are adsorbed at the carbidesremaining in the austenite, the carbides and austenite cover the grainboundaries of the ferrite, and, finally, it becomes possible to form astructure with a large amount of carbides present at the ferrite grainboundaries. For this reason, it is clear that the structure of thepresent invention cannot be formed by just simple annealing.

Below, the specific annealing conditions will be explained.

First Stage Annealing

Heating rate up to annealing temperature: 30° C./hour to 150° C./hour

Annealing temperature: 650° C. to 720° C.

Holding time at annealing temperature: 3 hours to 60 hours

The heating rate up to the first stage annealing temperature is made 30°C./hour to 150° C./hour. If the heating rate is less than 30° C./hour,time is required for raising the temperature and the productivity falls,so the heating rate is made 30° C./hour or more. Preferably, it is 40°C./hour or more.

On the other hand, if the heating rate is over 150° C./hour, thetemperature difference between the outer circumferential part and theinside part of the coil increases, scratches and seizing occur due tothe difference in heat expansion, and relief shapes are formed at thesteel plate surface. At the time of cold forging, cracks occur startingfrom the relief shapes and invite a drop in cold forgeability and a dropin impact resistance characteristic after carburizing, quenching, andtempering, so the heating rate is made 150° C./hour or less. Preferably,it is made 120° C./hour or less.

The annealing temperature in the first stage annealing (first stageannealing temperature) is made 650° C. to 720° C. If the first stageannealing temperature is less than 650° C., the carbides becomesinsufficient in stability and it becomes difficult to form carbidesremaining in the austenite in the second stage annealing, so the firststage annealing temperature is made 650° C. or more. Preferably it is670° C. or more.

On the other hand, if the annealing temperature exceeds 720° C., beforethe carbides rise in stability, austenite is formed and it becomesimpossible to control the above-mentioned changes in structure, so theannealing temperature is made 720° C. or less. Preferably it is 700° C.or less.

The holding time in the first stage annealing (first stage holding time)is made 3 hours to 60 hours. If the first stage holding time is lessthan 3 hours, the carbides become insufficient in stability and itbecomes difficult to form carbides remaining in the second stageannealing, so the first stage holding time is made 3 hours or more.Preferably it is 10 hours or more.

On the other hand, if the first stage holding time exceeds 60 hours, nofurther improvement in stability of the carbides can be expected.Furthermore, a drop in productivity is invited, so the first stageholding time is made 60 hours or less. Preferably it is 50 hours orless.

Second Stage Annealing

Heating rate up to annealing temperature: 1° C./hour to 80° C./hour

Annealing temperature: 725° C. to 790° C.

Holding time at annealing temperature: 3 hours to 50 hours

After finishing being held at the first stage annealing, the hot rolledsteel plate is heated up to the annealing temperature by a heating rateof 1° C./hour to 80° C./hour. If cooling without performing this secondstage annealing, the ferrite grain size does not become larger and theideal structure cannot be obtained.

In the second stage annealing, austenite is produced and grows from theferrite grain boundaries. By slowing the heating rate, it is possible tosuppress formation of nuclei of austenite. In the structure obtainedafter gradual cooling, it becomes possible to raise the rate of coverageof the grain boundaries of the carbides. For this reason, the heatingrate at the second stage annealing is preferably small.

If the heating rate is less than 1° C./hour, time is required forraising the temperature and the productivity falls, so the heating rateis made 1° C./hour or more. Preferably it is 10° C./hour or more.

On the other hand, if the heating rate exceeds 80° C./hour, thetemperature difference between the outer circumferential part and insidepart of the coil increases. Due to the large difference in heatexpansion due to deformation, scratches and seizing occur and reliefshapes are formed at the surface of the steel plate. At the time of coldforging, cracks form starting from the relief shapes leading to a dropin cold forgeability and a drop in impact resistance characteristicafter carburizing, quenching, and tempering, so the heating rate is made80° C./hour or less.

The annealing temperature in the second stage annealing (second stageannealing temperature) is made 725° C. to 790° C. If the second stageannealing temperature is less than 725° C., the amount of production ofaustenite becomes smaller. After cooling after the second stageannealing, the number ratio of carbides at the ferrite grain boundariesfalls and, further, the ferrite grain size becomes smaller. For thisreason, the second stage annealing temperature is made 725° C. or more.Preferably it is 735° C. or more.

On the other hand, if the second stage annealing temperature exceeds790° C., it becomes difficult to form carbides remaining in theaustenite and control to the above-mentioned change of structure becomesdifficult, so the second stage annealing temperature is made 790° C. orless. Preferably it is 780° C. or less.

The holding time in the second stage annealing (second stage holdingtime) is made 1 hour to 50 hours. If the second stage holding time isless than 1 hour, the amount of austenite which is produced is small,the carbides in the ferrite grains are not sufficiently dissolved, itbecomes difficult to increase the number ratio of carbides at theferrite grain boundaries, and, further, the ferrite grains becomesmaller in size, so the second stage holding time is made 1 hour ormore. Preferably, it is 5 hours or more.

On the other hand, if the second stage holding time exceeds 50 hours, itis difficult to make carbides remain in the austenite, so the secondstage holding time is made 50 hours or less. Preferably it is 45 hours.

Cooling After Annealing

Cooling stop temperature: 650° C.

Cooling rate: 1° C./hour to 100° C./hour

After finishing being held at the second stage annealing, the annealedhot rolled steel plate is gradually cooled down to 650° C. by a coolingrate of 1° C./hour to 100° C./hour. If the stop temperature of gradualcooling exceeds 650° C., due to the cooling rate subsequently exceeding100° C./hour down to room temperature, nontransformed austenitetransforms to pearlite or bainite, the hardness increases, and the coldforgeability falls, so the cooling stop temperature is made 650° C.

To cool the austenite formed in the second stage annealing and transformit to ferrite and to make carbon be adsorbed at the carbides remainingin the austenite, a slower cooling rate is preferable. If the coolingrate is less than 1° C./hour, the time required for cooling increasesand the productivity falls, so the cooling rate is made 1° C./hour ormore. Preferably it is 10° C./hour or more.

On the other hand, if the cooling rate exceeds 100° C./hour, austenitetransforms to pearlite and the steel plate increases in hardness so adrop in cold forgeability and a drop in impact resistancecharacteristics after carburizing, quenching, and tempering are invited,so the cooling rate is made 100° C./hour or less. Preferably it is 90°C./hour.

Here, the cooling stop temperature is the temperature where the coolingrate should be used for control. If cooling down to 650° C. by a coolingrate of 1° C./hour to 100° C./hour, the cooling down to 650° C. or lessis not particularly limited.

Note that, the atmosphere of the annealing is not limited to anyspecific atmosphere. For example, it may be any of an atmosphere of 95%or more of nitrogen, an atmosphere of 95% or more of hydrogen, and anair atmosphere.

As explained above, according to the method of the present invention ofintegrally managing the hot rolling conditions and annealing conditionsand controlling the structure of the steel plate, it is possible toproduce low carbon steel plate exhibiting excellent cold forgeability incold forging combining drawing and thickening and, furthermore,excellent in impact resistance characteristics after carburizing,quenching, and tempering.

EXAMPLES

Next, examples will be explained, but the levels in the examples areillustrations of conditions employed for confirming the workability andeffects of the present invention. The present invention is not limitedto these illustrations of conditions. The present invention can employvarious conditions so long as not deviating from the gist of the presentinvention and achieving the object of the present invention.

A continuously cast slab (steel ingot) having a chemical compositionshown in Table 1 was heated at 1240° C. for 1.8 hours, then was used forhot rolling. The finish hot rolling was ended at 890° C., the steel wascooled on a ROT by a 45° C./sec cooling rate down to 520° C. and wascoiled up at 510° C. to produce a hot rolled coil with a plate thicknessof 5.2 mm.

TABLE 1 Chemical composition (mass %) No. C Si Mn P S Al Cr Mo N O Ti BRemarks A 0.12 0.07 0.85 0.0154 0.0084 0.031 0.527 0.636 0.0068 0.00030.0095 0.0004 Invention steel B 0.13 0.03 0.76 0.0069 0.0046 0.011 1.4830.017 0.0057 0.0122 0.0039 0.0001 Invention steel C 0.17 0.18 0.340.0027 0.0019 0.004 0.996 0.204 0.0002 0.0030 0.0004 0.0000 Inventionsteel D 0.21 0.17 0.81 0.0133 0.0073 0.032 0.563 0.173 0.0086 0.01660.0030 0.0002 Invention steel E 0.23 0.24 0.64 0.0169 0.0095 0.046 1.9340.731 0.0094 0.0063 0.0043 0.0003 Invention Steel F 0.25 0.06 0.840.0189 0.0099 0.062 1.043 0.195 0.0048 0.0196 0.0033 0.0002 InventionSteel G 0.26 0.19 0.45 0.0112 0.0075 0.017 1.024 0.003 0.0053 0.01760.0088 0.0002 Invention Steel H 0.30 0.05 0.95 0.0100 0.0012 0.091 0.5860.591 0.0034 0.0170 0.0058 0.0004 Invention Steel I 0.34 0.28 0.920.0043 0.0017 0.078 1.905 0.860 0.0086 0.0188 0.0013 0.0000 InventionSteel J 0.36 0.01 0.51 0.0162 0.0006 0.069 0.705 0.943 0.0016 0.01250.0093 0.0002 Invention Steel K 0.39 0.10 0.85 0.0013 0.0074 0.033 1.6350.736 0.0131 0.0143 0.0049 0.0002 Invention Steel L 0.07 0.21 0.760.0166 0.0005 0.013 0.609 0.842 0.0050 0.0163 0.0086 0.0004 ComparativeSteel M 0.11 0.19 0.78 0.0211 0.0069 0.002 1.379 0.941 0.0183 0.01860.0087 0.0003 Comparative Steel N 0.14 0.24 0.79 0.0169 0.0020 0.0401.328 1.071 0.0174 0.0155 0.0063 0.0002 Comparative Steel O 0.15 0.130.58 0.0018 0.0098 0.031 0.449 0.291 0.0181 0.0171 0.0024 0.0003Comparative Steel P 0.20 0.19 0.31 0.0025 0.0090 0.108 0.525 0.7620.0195 0.0172 0.0084 0.0003 Comparative Steel Q 0.24 0.19 1.09 0.01960.0081 0.057 0.774 0.066 0.0066 0.0017 0.0039 0.0003 Comparative Steel R0.25 0.25 0.90 0.0099 0.0109 0.100 0.849 0.652 0.0178 0.0067 0.00170.0003 Comparative Steel S 0.27 0.31 0.36 0.0050 0.0063 0.094 0.8220.006 0.0019 0.0099 0.0019 0.0003 Comparative Steel T 0.29 0.26 0.630.0120 0.0049 0.003 2.236 0.011 0.0153 0.0092 0.0073 0.0001 ComparativeSteel U 0.45 0.19 0.64 0.0009 0.0025 0.072 1.150 0.008 0.0109 0.01540.0063 0.0004 Comparative Steel V 0.18 0.16 0.28 0.0129 0.0100 0.0480.917 0.961 0.0085 0.0001 0.0005 0.0002 Comparative Steel W 0.15 0.210.89 0.0076 0.0009 0.091 1.293 0.005 0.0058 0.0209 0.0042 0.0000Comparative Steel X 0.25 0.26 0.46 0.0070 0.0052 0.029 1.089 0.6100.0162 0.0161 0.0105 0.0001 Comparative Steel Y 0.28 0.20 0.76 0.01120.0026 0.083 1.193 0.713 0.0171 0.0083 0.0030 0.0006 Comparative Steel Z0.25 0.14 0.60 0.0033 0.0004 0.047 0.744 0.176 0.0208 0.0035 0.00540.0004 Comparative Steel

The hot rolled coil was pickled, the coil was loaded into a box-typeannealing furnace, the atmosphere was controlled to 95% hydrogen-5%nitrogen, then the coil was heated from room temperature up to 705° C.by a heating rate of 100° C./hour, was held at 705° C. for 36 hours tomake the temperature distribution inside the coil uniform, then washeated by a 5° C./hour heating rate up to 760° C., and, furthermore, washeld at 760° C. for 10 hours, then was cooled down to 650° C. by a 10°C./hour cooling rate, then was furnace cooled down to room temperatureto prepare a sample for evaluation of the characteristics.

The structure of the sample was observed by the above-mentioned method.The crack length at the sample after cold forging was measured by theabove-mentioned method.

The carburization of the thickened sample was performed by gascarburization. To make the carbon disperse from the inside of thefurnace atmosphere gas through the surface layer of the sample to theinside of the steel, the sample was treated by holding it at 940° C. for120 minutes inside a furnace controlled to a carbon potential of 0.5mass % C, then was furnace cooled down to room temperature.

Next, the sample was heated from room temperature to 840° C., then washeld for 20 minutes and quenched in 60° C. oil. The hardened sample washeld at 170° C. for 60 minutes, then air cooled for tempering to preparea carburized, quenched, and tempered sample.

The carburized, quenched, and tempered sample was evaluated for impactresistance by a drop weight test. FIG. 2 schematically shows an outlineof the drop weight test for evaluating the impact resistancecharacteristic of a carburized, quenched, and tempered sample. Thebottom of the cup of a carburized, quenched, and tempered cup-shapedsample 4 was fastened by a fixture. On a side surface of the cup, aweight 2 kg dropping weight (top side width: 50 mm, bottom side width:10 mm, height: 80 mm, and length: 110 mm) was allowed to freely dropfrom 4 m above to give an approximately 80J impact on the vertical wallpart of the sample 4. The sample was examined for the presence of anycracking and was evaluated for the impact resistance characteristic.

A sample with no fracture or breakage observed as a result of freedropping was evaluated as excellent in impact resistancecharacteristics, that is, “OK”, while a sample with a fracture orbreakage observed was evaluated as inferior in impact resistance, thatis, “NG”.

Table 2 shows the results of measurement and results of evaluation ofthe carbide size, pearlite area ratio, ferrite grain size, Vickershardness, ratio of the number of carbides at the ferrite grainboundaries to number of carbides in the ferrite grains, ratio of maximumcrack length to plate thickness at the vertical wall parts, and impactresistance characteristic in the prepared samples.

TABLE 2 No. of Ratio carbides of Ferrite Pearlite at grain maximumCarbide grain area Vickers boundaries/No. crack Impact size size ratehardness of carbides length resistance (μm) (μm) (%) (HV) inside grains(%) characteristic Remarks A-1 1.11 23 1.2 106.0 8.67 1.5 OK InventionSteel B-1 0.92 18.5 0.9 106.1 5.72 2.0 OK Invention Steel C-1 0.87 23.51.3 107.4 4.87 2.3 OK Invention Steel D-1 1.07 19.4 1.3 117.4 7.65 2.3OK Invention Steel E-1 0.78 15 1.3 126.7 1.08 3.8 OK Invention Steel F-10.99 17.2 0.7 115.7 6.75 2.4 OK Invention Steel G-1 0.88 19.5 0.3 117.54.90 2.9 OK Invention Steel H-1 1.09 17.7 0.5 118.0 8.44 2.2 OKInvention Steel I-1 0.84 13.1 0.5 140.2 1.76 3.4 OK Invention Steel J-10.98 18.3 1.1 114.7 6.68 2.3 OK Invention Steel K-1 0.9 14.9 0.7 127.35.63 3.3 OK Invention Steel L-1 1.08 26.2 0.6 92.2 9.29 15.5 NGComparative Steel M-1 0.96 19.7 1.1 114.4 0.12 15.4 NG Comparative SteelN-1 0.96 18.1 8.8 191.7 0.18 22.9 NG Comparative Steel O-1 1.06 22.7 0.1107.4 7.16 1.8 NG Comparative Steel P-1 0.97 25.2 0.6 110.0 0.61 13.7 NGComparative Steel Q-1 1.08 18.1 0.9 123.9 8.00 2.6 NG Comparative SteelR-1 1.03 17 1.2 128.8 7.06 18.8 NG Comparative Steel S-1 0.85 21.4 1.0124.3 4.23 13.6 NG Comparative Steel T-1 0.69 14.1 13.1 232.5 0.24 26.2NG Comparative Steel U-1 0.89 15.4 1.5 135.0 21.15 3.2 NG ComparativeSteel V-1 0.9 25.8 0.3 105.3 0.67 11.5 NG Comparative Steel W-1 0.9417.1 0.7 122 0.82 14.3 NG Comparative Steel X-1 0.89 19.3 0.1 121.8 0.5115.3 NG Comparative Steel Y-1 0.95 16.1 0.4 126.8 6.06 15.4 NGComparative Steel Z-1 1 19.1 0.8 116.1 0.68 14.2 NG Comparative Steel

As shown in Table 2, in Invention Steels A-1, B-1, C-1, D-1, E-1, F-1,G-1, H-1, I-1, J-1, and K-1, in each case, the ratio of the number ofcarbides at the ferrite grain boundaries to the number of carbides inthe ferrite grains is over 1, the Vickers hardness is 100 HV to 180 HV,and the cold forgeability and impact resistance characteristics aftercarburizing, quenching, and tempering are excellent.

As opposed to this, in Comparative Steel L-1, the amount of C is low andthe hardness before cold forging is less than 100 HV, so the coldforgeability is low. In Comparative Steels M-1, P-1, and Z-1, P, Al, andN are excessively contained and, at the second stage annealing, theamount of segregation at the y/a interfaces is large, so formation ofcarbides at the grain boundaries is suppressed.

In Comparative Steel S-1, Si is excessively contained and ductility ofthe ferrite is low, so the cold forgeability is low. In ComparativeSteels N-1 and T-1, Mo and Cr are excessively contained, so carbidesfinely disperse inside the ferrite grains and the hardness exceeds 180HV. In Comparative Steel Q-1, Mn is excessively contained, so the impactresistance characteristic after carburizing, quenching, and tempering isremarkably low.

In Comparative Steel 0-1, the amount of Cr is small and the austenitegrains at the surface layer abnormally coarsen at the time ofcarburization, so the impact resistance is low. In Comparative SteelR-1, S is excessively contained, so coarse MnS is formed in the steeland the cold forgeability is low. In Comparative Steel U-1, C isexcessively contained, so coarse carbides form inside the thickenedlayer of the steel and coarse carbides remain even after the carburizingand quenching, so the impact resistance characteristic is low.

In Comparative Steel V-1, the amount of Mn is small and the carbides arehard to raise in stability, so the cold forgeability and impactresistance characteristic after carburizing, quenching, and temperingare low. In Comparative Steels W-1 and X-1, 0 and Ti are excessivelycontained, so the oxides and TiC present in the ferrite grains becomesite for formation of carbides in gradual cooling after dual phaseregion annealing, the formation of carbides at the grain boundaries issuppressed, and the cold forgeability is low. In Comparative Steel Y-1,B is excessively contained, so the cold forgeability is low.

Next, to investigate the effects of the manufacturing conditions, slabshaving the chemical compositions A, B, C, D, E, F, G, H, I, J, and Kshown in Table 1 were hot rolled and annealed under the conditions shownin Table 3 to prepare annealed samples of hot rolled plates of athickness of 5.2 mm.

Table 4 shows the results of measurement and results of evaluation ofthe carbide size, pearlite area ratio, ferrite grain size, Vickershardness, ratio of the number of carbides at the ferrite grainboundaries to number of carbides in the ferrite grains, ratio of maximumcrack length to plate thickness at the vertical wall parts, and impactresistance characteristics in the prepared samples.

TABLE 3 Hot rolling conditions Annealing conditions Finish ROT 1st stage2nd stage Cooling hot cooling Heating Heating rate Heating Soakingrolling rate Coiling rate Holding Holding rate Holding Holding down totemp. time temp. (° C./ temp. (° C./ temp. time (° C./ temp. time 650°C. (° C.) (hours) (° C.) sec) (° C.) hour) (° C.) (hours) hour) (° C.)(hours) (° C./hour) Remarks A-2 — — 837.1 63 449.2 143.7 653.1 44.1 48.1745.2 49.7 26.8 Inv. steel B-2 1162 0.5 905 35 409.3 101.3 706.5 32.92.2 773.5 13.6 84.6 Inv. steel C-2 1196 2.5 882.3 64 388 67.9 695.6 36.519.5 735.7 20.7 33.8 Comp. steel D-2 1060 5.6 826 21 496.8 55.9 683.249.2 31.4 789 41.4 99.5 Inv. steel E-2 1241 3.4 639 33 526.7 84.5 687.543 40.5 775.1 4.6 72.4 Comp. steel F-2 1271 0.8 714.7 91 412.1 71.4709.2 63.1 23.8 769.7 14.4 10.8 Comp. steel G-2 1265 5.6 844.4 52 627133.6 671.7 44.2 28.9 750 39.7 92.1 Comp. steel H-2 1287 5.9 752.9 30567.5 139.4 709 22 16.4 731.9 28.9 57.7 Inv. steel I-2 1030 5.6 850.3100 541.1 44.5 695.4 51.9 29.9 788.7 30.5 76.7 Inv. steel J-2 1258 0.6741.9 76 543.9 35.9 720 33.7 78.8 773 1.3 1.4 Inv. steel K-2 1152 5.9703 41 516.2 44.8 681.1 46.5 2.8 721 35.3 72 Comp. steel A-3 1138 4.5824.7 53 472.1 85.9 687.2 35.8 88 773 29.2 92.8 Comp. steel B-3 1254 1.7717.5 70 526.9 42.6 715.1 15.6 48 740.3 1.8 6.7 Comp. steel C-3 1292 2.5870.5 35 535.8 36.2 669 47.7 37.9 784.7 15.9 95 Inv. steel D-3 1010 1.4812.3 69 403.5 124.9 709.2 11 72.3 734.5 8.8 0.4 Comp. steel E-3 10070.8 689.5 93 466.5 158 700 54.6 55.8 775.8 45.6 16.7 Comp. steel F-31062 3.6 656.1 25 567.7 126.5 699.9 42.7 24.9 731.2 29.5 13.1 Inv. steelG-3 1137 2.8 851.7 48 446.5 94.4 682.2 14.8 77.1 757.8 42.6 16.4 Inv.steel H-3 1089 5.6 767.6 10 468.4 146.7 685.1 49.1 8 768.9 45.6 34.8Inv. steel I-3 1288 1.7 682.3 55 443 109.2 683.1 39.2 63.7 769.9 29.9116 Comp. steel J-3 1037 5.6 879 53 519.9 59.1 719.2 59.8 77.5 727 5747.3 Comp. steel K-3 — — 664.6 85 407.9 75.7 665.4 32 40.2 777 21.4 32Inv. steel A-4 1164 2.5 710.6 80 503.2 135.8 669.3 47.1 71.3 731.3 25.869.3 Inv. steel B-4 1136 4.8 695.9 28 445 90.1 701.4 21 0.2 776.1 42.987.4 Comp. steel C-4 1046 1.1 667.3 80 405.1 54.9 710.4 47.8 56.4 80417.1 97.5 Comp. steel D-4 1253 4.5 708.1 29 493.9 74 737 57.3 38.4 755.733.6 55.4 Comp. steel E-4 1018 5.0 870.2 54 560.4 141.7 710.9 9.2 19.7748 16.2 96.5 Inv. steel F-4 1229 6.2 962 71 475.1 99.3 672.1 24 28.4736.3 24.1 7.9 Comp. steel G-4 1139 6.4 749.8 19 535 144.8 639 39.3 18783.9 32.5 91.3 Comp. steel H-4 1119 4.8 731.7 66 456.1 13 668.3 46.655.8 744.7 19 43.8 Comp. steel I-4 1101 4.2 686 33 569.4 73.6 656.1 31.120.3 740.4 31.5 24.2 Inv. steel J-4 1042 0.6 750.8 69 545.5 40 719.4 1.471.9 729.1 38.3 11.9 Comp. steel K-4 1268 4.5 912.2 44 521.5 92.5 652.254.7 39.2 745.5 37.4 48.1 Inv. steel

TABLE 4 No. of Ratio carbides of Ferrite Pearlite at grain maximumCarbide grain area Vickers boundaries/No. crack Impact size size ratehardness of carbides length resistance (μm) (μm) (%) (HV) inside grains(%) characteristic Remarks A-2 0.71 19.6 1.2 103.5 9.19 1.3 OK InventionSteel B-2 0.73 22.8 0.7 101.6 1.90 2.9 OK Invention Steel C-2 1 18.4 1.0111.4 6.83 2.1 NG Comparative Steel D-2 0.57 25.0 1.4 106.5 2.16 3.3 OKInvention Steel E-2 0.54 13.7 0.9 126.6 2.57 4.3 OK Comparative SteelF-2 2.23 38.7 0.8 98.4 4.24 12.9 NG Comparative Steel G-2 0.28 16.2 1.1120.0 1.56 14.9 NG Comparative Steel H-2 0.55 10.2 0.4 127.3 1.69 5.2 OKInvention Steel I-2 0.57 18.4 0.2 125.5 1.89 4.8 OK Invention Steel J-21.95 29.5 1.1 104.4 18.06 0.7 OK Invention Steel K-2 0.76 8.8 0.7 136.30.08 17.9 NG Comparative Steel A-3 0.59 23.5 0.1 97.9 2.45 22.6 NGComparative Steel B-3 0.7 8.3 0.1 129.3 0.16 16.9 NG Comparative SteelC-3 0.48 25.6 1.1 104.3 2.71 2.8 OK Invention Steel D-3 2.34 16.9 0.1128.6 21.84 21.5 NG Comparative Steel E-3 0.86 28.4 0.6 107.9 3.32 5.0NG Comparative Steel F-3 0.58 10.3 0.2 124.9 1.71 5.0 OK Invention SteelG-3 0.72 25.5 1.4 107.8 7.56 1.8 OK Invention Steel H-3 0.78 22.1 1.4105.2 3.76 2.4 OK Invention Steel I-3 0.49 16.6 9.5 186.2 2.05 20.5 NGComparative Steel J-3 1.18 17.9 9.6 189.7 1.52 14.3 NG Comparative SteelK-3 0.74 19.1 0.5 113.9 9.88 1.8 OK Invention Steel A-4 0.47 11.1 0.3119.0 2.75 3.7 OK Invention Steel B-4 0.57 26.2 0.8 93.1 2.08 2.5 OKComparative Steel C-4 0.84 40.2 10.6 195.0 1.59 19.2 NG ComparativeSteel D-4 1.26 26.4 8.9 196.5 1.23 14.5 NG Comparative Steel E-4 0.3711.2 0.7 131.7 1.29 6.2 OK Invention Steel F-4 0.7 11.6 1.4 121.7 13.916.2 NG Comparative Steel G-4 0.48 24.6 12.4 211.2 3.03 18.4 NGComparative Steel H-4 0.62 12.6 0.3 120.9 5.65 2.9 OK Comparative SteelI-4 0.25 9.7 0.0 143.3 2.06 5.7 OK Invention Steel J-4 0.34 13.1 14.2218.2 3.13 18.7 NG Comparative Steel K-4 0.38 12.4 0.1 125.8 3.44 3.8 OKInvention Steel

In Comparative Steel E-3, the finish hot rolling temperature is low, therolling load increases, and the productivity falls. In Comparative SteelD-2, the finish hot rolling temperature is high and scale flaws form atthe surface of the steel plate, so when subjected to a wear resistancetest after quenching and tempering, fractures and peeling occur startingfrom the scale flaws and the wear resistance characteristic falls. InComparative Steel F-2, the cooling rate at the ROT (Run Out Table) isslow and a drop of productivity and formation of scale flaws areinvited.

In Comparative Steel C-4, the cooling rate at the ROT is 100° C./sec andthe outermost layer part of the steel plate is excessively cooled, sofine cracks formed at the outermost layer part. In Comparative SteelC-2, the coiling temperature is low, large amounts of bainite,martensite, and other low temperature transformed structures are formedcausing embrittlement, fractures frequently form at the time of pay outfrom the hot rolled coil, and the productivity falls. Furthermore, in asample taken from a cracked piece, the wear resistance characteristic islow.

In Comparative Steel G-2, the coiling temperature is high, bulkypearlite of lamellar spacing is formed in the hot rolled structure, theneedle shaped coarse carbides become high in thermal stability, and theabove carbides remain in the steel plate even after two-stage step typeannealing, so the machinability is low. In Comparative Steel H-4, theheating rate in the first stage annealing of the two-stage step typeannealing is slow, so the productivity is low.

In Comparative Steel E-3, the heating rate in the first stage annealingis fast, so the temperature difference between the inside part andinside and outside circumferential parts of the coil becomes larger,scratches and seizing occur due to the difference in thermal expansion,and, when used for evaluating and testing the wear resistancecharacteristic after quenching and tempering, cracks and peeling occurfrom the flaw parts and the wear resistance characteristic falls.

In Comparative Steel G-4, the holding temperature in the first stageannealing (annealing temperature) is low, the coarsening treatment ofcarbides at the Ac1 point or less is insufficient, and the carbides areinsufficient in thermal stability, so carbides remaining at the secondstage annealing are reduced and pearlite transformation in the structureafter gradual cooling cannot be suppressed, so the machinability is low.

In Comparative Steel D-4, the holding temperature in the first stageannealing (annealing temperature) is high, austenite is formed duringthe annealing, and the carbides cannot be raised in stability, sopearlite is formed after annealing, the Vickers hardness exceeds 180 HV,and the machinability is low. In Comparative Steel J-4, the holding timein the first stage annealing is short and the stability of carbidescannot be raised, so the machinability is low.

In Comparative Steel F-2, the holding time in the first stage annealingis long, the productivity is low, and further seizing flaws occur andthe wear resistance characteristic is low. In Comparative Steel B-4, theheating rate in the second stage annealing in the two-stage step typeannealing is slow, so the productivity is low. In Comparative Steel A-3,the heating rate at the second stage annealing is fast, so thetemperature difference between the inside part and outer circumferentialpart of the coil become greater, scratches and seizing occur due to thelarge difference in heat expansion due to deformation, and the wearresistance characteristic after quenching and tempering is low.

In Comparative Steel K-2, the holding temperature in the second stageannealing (annealing temperature) is low, the amount of production ofaustenite is small, and the ratio of number of carbides at the ferritegrain boundaries cannot be increased, so the machinability is low. InComparative Steel C-4, the holding temperature at the second stageannealing (annealing temperature) is high and dissolution of thecarbides during the annealing is promoted, so it becomes difficult toform carbides at the grain boundaries after the gradual cooling andfurther pearlite is produced, the Vickers hardness exceeds 180 HV, andthe machinability is low.

In Comparative Steel J-3, the holding time at the second stage annealingis long and dissolution of the carbides is promoted, so themachinability is low. In Comparative Steel D-3, the cooling rate fromsecond stage annealing to 650° C. is slow, the productivity is low,coarse carbides are formed in the structure after gradual cooling,cracks are formed starting from the coarse carbides at the time of coldforging, and the cold forgeability falls. In Comparative Steel 1-3, thecooling rate from second stage annealing to 650° C. is fast, thepearlite transformation occurs at the time of cooling, and the hardnessincreases, so the cold forgeability is low.

Next, to investigate the allowable contents of the other elements,continuously cast slabs (steel ingots) having the chemical compositionsshown in Table 5 and Table 6 (continuation of Table 5) were heated at1240° C. for 1.8 hours, then were used for hot rolling. The finish hotrolling was ended at 890° C., the steels were cooled on a ROT by a 45°C./sec cooling rate down to 520° C. and were coiled up at 510° C. toproduce hot rolled coils with a plate thickness of 5.2 mm.

TABLE 5 Chemical composition (mass %) C Si Mn P S Al N O Ti Cr Mo B Nb VAA 0.13 0.01 0.96 0.0076 0.0063 0.011 0.0139 0.0112 0.0043 0.509 0.8690.0001 0.028 AB 0.16 0.25 0.70 0.0063 0.0087 0.083 0.0162 0.0119 0.00280.618 0.680 0.0001 0.029 AC 0.18 0.11 0.97 0.0007 0.0043 0.073 0.00670.0076 0.0095 1.199 0.402 0.0004 AD 0.22 0.29 0.69 0.0145 0.0020 0.0070.0077 0.0005 0.0015 1.400 0.807 0.0003 0.012 AE 0.22 0.21 0.61 0.00670.0072 0.002 0.0008 0.0093 0.0072 1.130 0.422 0.0002 0.004 AF 0.27 0.160.60 0.0098 0.0032 0.079 0.0011 0.0011 0.0042 1.197 0.010 0.0003 0.070AG 0.28 0.04 0.79 0.0075 0.0035 0.049 0.0116 0.0137 0.0028 0.862 0.8020.0001 0.040 AH 0.28 0.07 0.43 0.0064 0.0058 0.006 0.0089 0.0025 0.00361.346 0.510 0.0000 0.046 0.065 AI 0.30 0.10 0.80 0.0047 0.0045 0.0610.0196 0.0027 0.0057 0.961 0.002 0.0002 0.022 0.094 AJ 0.31 0.27 0.600.0121 0.0046 0.077 0.0105 0.0153 0.0029 1.649 0.013 0.0001 0.002 0.084AK 0.36 0.07 0.60 0.0093 0.0040 0.028 0.0016 0.0022 0.0043 1.722 0.0090.0004 AL 0.36 0.29 0.89 0.0187 0.0055 0.048 0.0190 0.0114 0.0031 0.8810.893 0.0004 0.092 AM 0.37 0.06 0.84 0.0002 0.0002 0.072 0.0102 0.00920.0020 0.581 0.345 0.0001 0.015 0.037 AN 0.37 0.19 0.45 0.0139 0.00490.033 0.0050 0.0147 0.0055 0.540 0.077 0.0001 0.039 0.012 AO 0.39 0.291.00 0.0176 0.0009 0.068 0.0142 0.0176 0.0070 1.951 0.330 0.0002 AP 0.390.27 0.82 0.0075 0.0025 0.025 0.0116 0.0076 0.0077 0.981 0.387 0.00020.062 AQ 0.39 0.17 0.43 0.0194 0.0034 0.014 0.0037 0.0036 0.0014 1.4800.626 0.0002 0.018 Chemical composition (mass %) Cu W Ta Ni Sn Sb As MgCa Y Zr La Ce Remarks AA 0.095 0.021 0.028 0.034 0.029 0.001 0.047 0.020Invention steel AB 0.082 0.086 0.053 0.086 0.016 0.024 0.043 0.005 0.0280.034 0.046 Invention steel AC 0.008 0.031 0.029 0.008 0.002 0.045 0.0110.039 0.006 Invention steel AD 0.019 0.019 0.053 0.013 0.029 0.034 0.026Invention steel AE 0.054 0.003 0.085 0.027 0.002 0.001 0.003 0.026 0.031Invention steel AF 0.038 0.014 0.048 0.006 0.042 0.039 0.009 0.013Invention steel AG 0.076 0.073 0.038 0.009 0.049 0.017 0.045 0.026Invention steel AH 0.076 0.007 0.002 0.024 0.019 0.012 0.029 0.011Invention steel AI 0.092 0.012 0.041 0.038 0.008 0.010 0.014 Inventionsteel AJ 0.086 0.080 0.048 0.027 0.041 0.021 0.030 0.001 Invention steelAK 0.003 Invention steel AL 0.095 0.042 0.078 0.005 0.006 0.046 0.032Invention steel AM 0.058 0.062 0.048 0.011 0.046 0.006 0.002 0.021 0.042Invention steel AN 0.067 0.093 0.021 0.006 0.044 0.019 0.016 0.001 0.040Invention steel AO 0.005 Invention steel AP Invention steel AQ 0.0680.002 0.005 0.029 0.033 0.031 0.004 0.023 Invention steel

TABLE 6 (Continuation of Table 5) Chemical composition (mass %) C Si MnP S Al N O Ti Cr Mo B Nb V AR 0.12 0.13 0.95 0.0195 0.0083 0.06 0.01320.0033 0.0064 0.554 0.014 0.0005 0.045 0.006 AS 0.15 0.11 0.52 0.01220.0050 0.008 0.0177 0.0011 0.0008 0.773 0.632 0.0001 0.002 0.038 AT 0.150.06 0.70 0.0115 0.0060 0.036 0.0139 0.0100 0.0083 0.663 1.046 0.00050.001 AU 0.15 0.09 0.57 0.0165 0.0029 0.048 0.0154 0.0040 0.0074 0.8110.314 0.0002 AV 0.16 0.18 0.54 0.0058 0.0064 0.086 0.0051 0.0079 0.00701.406 0.161 0.0002 0.116 AW 0.19 0.27 0.67 0.0199 0.0009 0.083 0.00040.0009 0.0054 0.513 0.728 0.0002 0.032 AX 0.19 0.24 1.00 0.0050 0.00100.084 0.0016 0.0162 0.0095 0.904 0.841 0.0004 0.080 0.077 AY 0.24 0.240.53 0.0094 0.0005 0.018 0.0060 0.0032 0.0084 1.688 0.811 0.0003 0.044AZ 0.24 0.03 0.42 0.0050 0.0018 0.088 0.0076 0.0008 0.0002 1.216 0.8440.0004 0.067 0.025 BA 0.28 0.08 0.58 0.0005 0.0021 0.045 0.0094 0.00470.0032 2.037 0.076 0.0000 0.050 BB 0.30 0.03 0.60 0.0044 0.0094 0.0640.0015 0.0021 0.0075 1.285 0.513 0.0002 0.065 BC 0.31 0.02 0.57 0.01300.0074 0.041 0.0050 0.0157 0.0001 1.379 0.004 0.0001 0.032 0.051 BD 0.330.30 0.84 0.0140 0.0095 0.057 0.0085 0.0169 0.0030 1.648 0.182 0.0000 BE0.34 0.29 0.58 0.0004 0.0087 0.072 0.0011 0.0092 0.0019 1.094 0.0070.0003 0.091 0.079 BF 0.34 0.29 0.94 0.0149 0.0024 0.022 0.0167 0.00360.0050 1.561 0.856 0.0003 0.097 BG 0.36 0.36 0.60 0.0157 0.0088 0.0860.0198 0.0064 0.0012 0.934 0.268 0.0001 BH 0.36 0.28 0.63 0.0099 0.00910.032 0.0098 0.0051 0.0029 1.624 0.011 0.0000 0.047 0.044 BI 0.37 0.141.17 0.0014 0.0048 0.094 0.0151 0.0113 0.0003 1.210 0.003 0.0000 0.085BJ 0.39 0.19 0.95 0.0071 0.0012 0.094 0.0014 0.0186 0.0082 1.317 0.8490.0002 0.051 0.118 BK 0.45 0.17 0.61 0.0115 0.0055 0.054 0.0024 0.01670.0086 0.688 0.118 0.0000 0.043 Chemical composition (mass %) Cu W Ta NiSn Sb As Mg Ca Y Zr La Ce Remarks AR 0.075 0.084 0.024 0.027 0.049 0.052Comparative steel AS 0.025 0.055 0.041 0.003 0.058 0.001 0.040 0.0440.013 0.026 Comparative steel AT 0.012 0.048 0.011 0.011 0.031 0.0260.029 Comparative steel AU 0.098 0.095 0.047 0.008 0.037 0.053 0.011Comparative steel AV 0.078 0.005 0.048 0.027 0.035 0.040 0.006Comparative steel AW 0.131 0.021 0.036 0.012 0.025 0.004 0.036 0.0400.015 0.006 Comparative steel AX 0.087 0.045 0.011 0.040 0.014 0.0430.058 0.020 0.041 0.046 Comparative steel AY 0.014 0.095 0.025 0.0050.052 0.027 0.040 0.027 0.027 Comparative steel AZ 0.097 0.073 0.1060.009 0.044 0.017 0.020 0.022 0.025 0.025 0.014 Comparative steel BA0.031 0.036 0.020 0.008 0.027 0.012 0.050 0.026 0.017 Comparative steelBB 0.051 0.091 0.072 0.005 0.055 0.019 0.012 0.002 Comparative steel BC0.034 0.003 0.105 0.064 0.005 0.005 0.029 0.028 0.049 Comparative steelBD 0.059 Comparative steel BE 0.099 0.019 0.049 0.038 0.043 0.062 0.0280.022 0.032 Comparative steel BF 0.023 0.049 0.079 0.049 0.001 0.0010.002 0.051 Comparative steel BG 0.070 0.005 0.040 0.047 0.016 0.0230.032 0.002 Comparative steel BH 0.107 0.014 0.068 0.045 0.003 0.0140.018 Comparative steel BI 0.008 0.028 0.017 0.029 0.006 0.009 0.0270.037 0.038 Comparative steel BJ 0.077 0.096 0.076 0.012 0.046 0.0360.047 0.011 0.022 0.043 0.011 Comparative steel BK 0.083 0.048 0.0160.025 0.039 0.004 0.046 Comparative steel

The hot rolled coils were pickled, the hot rolled coils were loaded intoa box-type annealing furnace, the atmosphere was controlled to 95%hydrogen-5% nitrogen, then the coils were heated from room temperatureup to 705° C. by a heating rate of 100° C./hour, were held at 705° C.for 36 hours to make the temperature distribution inside the coilsuniform, then were heated by a 5° C./hour heating rate up to 760° C.,and, furthermore, were held at 760° C. for 10 hours, then were cooleddown to 650° C. by a 10° C./hour cooling rate, then were furnace cooleddown to room temperature to prepare samples for evaluation of thecharacteristics.

Note that, the structures of the samples were observed by theabove-mentioned method while the crack lengths present in the samplesafter cold forging were measured by the above-mentioned method.

Table 7 shows the results of measurement and results of evaluation ofthe carbide size, pearlite area ratio, ferrite grain size, Vickershardness, ratio of the number of carbides at the ferrite grainboundaries to number of carbides in the ferrite grains, ratio of maximumcrack length to plate thickness at the vertical wall parts, and impactresistance characteristics in the prepared samples.

TABLE 7 No. of carbides at t grain Ratio boundaries/ of Ferrite PearliteNo. of maximum Carbide grain area Vickers carbides crack Impact sizesize rate hardness inside length resistance (μm) (μm) (%) (HV) grains(%) characteristic Remarks AA-1 1.14 23.5 1.2 103.1 9.53 0.8 OKInvention steel AB-1 1.04 20 0.7 119.7 7.19 1.5 OK Invention steel AC-10.99 17.5 0.7 117.2 7.03 1.4 OK Invention steel AD-1 0.89 16.4 0.3 127.05.25 2.0 OK Invention steel AE-1 0.94 17.8 0.7 119.5 5.81 1.7 OKInvention steel AF-1 0.88 16.8 1.4 121.6 4.71 1.9 OK Invention steelAG-1 1.02 17.6 0.0 114.8 7.35 1.3 OK Invention steel AH-1 0.86 19.1 0.7111.6 4.56 1.6 OK Invention steel AI-1 1 17.4 0.0 119.7 6.75 1.6 OKInvention steel AJ-1 0.79 15.3 0.8 132.5 3.60 2.5 OK Invention steelAK-1 0.79 15.2 1.4 122.3 3.66 2.1 OK Invention steel AL-1 1.03 16.1 0.8136.0 7.59 2.0 OK Invention steel AM-1 1.06 17.6 0.3 120.6 7.63 1.5 OKInvention steel AN-1 0.97 19.1 0.1 124.2 5.46 1.9 OK Invention steelAO-1 0.83 13.4 1.0 142.6 4.61 2.6 OK Invention steel AP-1 0.98 16 0.2135.4 6.32 2.2 OK Invention steel AQ-1 0.81 17.3 1.1 126.3 3.88 2.2 OKInvention steel AR-1 1.1 22.8 0.7 111.3 0.15 14.9  NG Comparative steelAS-1 0.99 22 1.0 105.6 0.28 14.0  NG Comparative steel AT-1 1.07 21.37.8 188.0 0.19 22.6  NG Comparative steel AU-1 0.98 21.2 1.3 106.4 6.8921.9  NG Comparative steel AV-1 0.83 18.1 13.1 243.1 0.44 26.4  NGComparative steel AW-1 0.94 19.2 0.6 122.5 0.04 16.5  NG Comparativesteel AX-1 1.06 17.5 0.4 126.0 10.11  21.4  NG Comparative steel AY-11.06 16.7 1.5 123.4 4.01 14.8  NG Comparative steel AZ-1 0.81 19.8 0.4107.8 0.89 12.4  NG Comparative steel BA-1 0.87 15.2 11.4 234.4 0.2526.3  NG Comparative steel BB-1 0.73 16.7 1.0 115.2 0.11 15.5  NGComparative steel BC-1 0.9 16.8 9.8 212.6 0.94 21.2  NG Comparativesteel BD-1 0.88 14.1 1.2 138.2 4.53 13.4  NG Comparative steel BE-1 0.916.4 0.6 133.4 4.83 11.2  NG Comparative steel BF-1 0.85 14.6 1.2 137.20.73 16.2  NG Comparative steel BG-1 0.9 16.5 0.7 138.8 4.97 13.8  NGComparative steel BH-1 0.93 14.9 10.6 228.7 0.85 23.4  NG Comparativesteel BI-1 0.91 16.6 0.3 128.4 7.24 3.0 NG Comparative steel BJ-1 0.8214.5 11.4 240.6 0.24 26.8  NG Comparative steel BK-1 1.01 16.6 0.8 131.522.08  1.1 NG Comparative steel

As shown in Table 7, in each of Invention Steels AA-1, AB-1, AC-1, AD-1,AE-1, AF-1, AG-1, AH-1, AI-1, AJ-1, AK-1, AL-1, AM-1, AN-1, AO-1, AP-1,and AQ-1, the ratio of the number of carbides at the ferrite grainboundaries to the number of carbides in the ferrite grains exceeds 1,the Vickers hardness is 100 HV to 180 HV, and the cold forgeability andimpact resistance characteristic after carburizing, quenching, andtempering are excellent.

As opposed to this, in each of Comparative Steels AR-1, AS-1, AW-1,AZ-1, BB-1, and BF-1, La, As, Cu, Ni, Sb, and Ce are excessivelycontained and the amount of segregation at the y/a interface becomesgreater at the time of second stage annealing, so formation of carbidesat the grain boundaries is suppressed. In Comparative Steel BG-1, Si isexcessively contained and the ductility of the ferrite is low, so thecold forgeability is low.

In each of Comparative Steels AT-1, AV-1, BA-1, BC-1, BH-1, and BJ-1,Mo, Nb, Cr, Ta, W, and V are respectively excessively contained, socarbides finely disperse inside the ferrite grains and the hardnessexceeds 180 HV. In Comparative Steel BF-1, Mn is excessively contained,so the impact resistance characteristic after carburizing, quenching,and tempering is remarkably low.

In each of Comparative Steels AU-1, AX-1, AY-1, and BE-1, Zr, Ca, Mg,and Y are respectively excessively contained, coarse oxides ornonmetallic inclusions are formed in the steel, cracks form startingfrom the coarse oxides or coarse nonmetallic inclusions at the time ofcold forging, and the cold forgeability falls. In Comparative SteelBD-1, Sn is excessively contained, the ferrite becomes brittle, and thecold forgeability is low. In Comparative Steel BK-1, C is excessivelycontained, so coarse carbides form at the inside of the increasedthickness part of the steel, coarse carbides remain even aftercarburizing and quenching, and the impact resistance characteristic alsofalls.

Next, to investigate the effects of the manufacturing conditions, slabshaving the chemical compositions of AA, AB, AC, AD, AE, AF, AG, AH, AI,AJ, AK, AL, AM, AN, AO, AP, and AQ shown in Table 5 were hot rolled andannealed under the conditions shown in Table 8 to fabricate annealedsamples of hot rolled plates of thicknesses of 5.2 mm.

TABLE 8 Hot rolling conditions Annealing conditions Finish ROT 1st stage2nd stage hot cooling Heating Heating Cooling rate Heating Soakingrolling rate Coiling rate Holding Holding rate Holding Holding down totemp. time temp. (° C./ temp. (° C./ temp. time (° C./ temp. time 650°C. (° C.) (hours) (° C.) sec) (° C.) hour) (° C.) (hours) hour) (° C.)(hours) (° C./hour) Remarks AA-2 1143 0.4 836.3 11 403.5 124 644 31.936.2 774 14.2 47.5 Comp. steel AB-2 1199 4.5 939.4 17 435.8 123.4 692.338.9 52.5 766.7 34.7 35.3 Inv. steel AC-2 1043 4.2 641 61 415.4 75.8703.9 50.4 24.2 741.7 15.1 88.1 Comp. steel AD-2 1272 0.8 812.8 65 467.496.7 692.7 59.9 39.6 796 7.3 60.5 Comp. steel AE-2 1164 0.5 902.3 53 554131.8 670.2 15.7 78.3 734.5 49.8 66.4 Inv. steel AF-2 1238 4.5 832 42435.3 40.5 671.4 2.7 29.3 757.4 11.8  6.1 Comp. steel AG-2 1226 3.4830.5 14 503.9 171 688.5 28.1 14.6 747 10.8 51   Comp. steel AH-2 10501.4 833 48 558.2 102 706.7 22.2 26.9 744.5 49.7 80.7 Inv. steel AI-21067 2.8 690.7 26 452 45.7 662.2 31.5 53.6 737.7 10.2 29.6 Inv. steelAJ-2 1008 1.4 850.8 53 564.3 70.8 693.9 29.9 4.9 787.8 41.6 83.3 Inv.steel AK-2 1166 4.8 729.7 23 510 43.9 667.2 13.6 32 753 7.5 40.4 Inv.steel AL-2 1150 1.7 802.3 60 516.2 95.6 660.4 47.1 21.8 742.7 30 36.1Inv. steel AM-2 1083 1.4 779.6 66 534.9 86.2 685.7 25.5 22.8 754 28.515.4 Inv. steel AN-2 1084 1.1 799.6 96 552.8 74.1 712 52.5 53.4 739 43.417.4 Inv. steel AO-2 1194 2.2 710.1 16 592.8 42.5 689.2 50.7 67.4 74827.1 86.6 Inv. steel AP-2 1101 2.5 800.4 61 480.4 119 702 12.1 71.1781.4 27.7 123    Comp. steel AQ-2 — — 842.1 24 403.5 43.2 693.2 29.723.7 773.9 31.1 48.3 Inv. steel AA-3 1173 3.4 758.8 45 521.1 81.2 686.818.4 8.6 751 21.3 27   Inv. steel AB-3 1023 2.8 834.7 10 500.4 114 677.327.6 59.8 745.4 21.5 64.7 Inv. steel AC-3 1296 0.6 725.4 87 416.2 142.7682.2 58.8 40.3 740.1 10.7  4.9 Inv. steel AD-3 1143 3.9 946.1 91 588.636.6 677.3 53.1 8.8 737.3 6.9 72   Inv. steel AE-3 1110 0.6 874.7 44454.5 106.6 678.1 31.5 57.6 770 18.1 42.7 Inv. steel AF-3 1054 5.9 684.697 526.2 49.1 693.9 32.5 44.5 777.1 28.3 40.4 Inv. steel AG-3 1163 4.2773 63 501.3 52.9 701.8 9.8 18.2 751.3 44.4 43.7 Inv. steel AH-3 11220.5 884.9 36 613 84.2 686.6 58.5 2.9 731.3 33.9 39   Comp. steel AI-31245 7.0 797.5 44 433.1 136.2 698.7 54.9 91 778 48.4 60.6 Comp. steelAJ-3 1281 4.8 895.7 100 562.9 127 673.4 18.9 13.2 748 28.7 78.6 Inv.steel AK-3 1164 5.0 910.4 86 486.1 123.9 653.8 56.4 4.4 758.9 2.1 41.5Comp. steel AL-3 1212 0.8 691 87 486.5 61.8 693.4 30.6 15.3 709 23.585.9 Comp. steel AM-3 1098 4.5 798.7 44 543.5 130.7 726 4.9 46 770.243.9 19.4 Comp. steel AN-3 1022 7.0 879.7 48 391 146.2 657 13.7 10.9765.1 32.7 89.9 Comp. steel AO-3 1094 3.9 725.5 67 488.4 137.2 671.442.5 71.1 778.9 19.5 40.3 Inv. steel AP-3 1083 1.1 919.3 99 573.9 42.3715.3 25.6 25.8 730.3 14.4 97.7 Inv. steel AQ-3 1096 3.1 743.6 22 527.4113.3 652.7 46.3 13.6 751.3 14.3   0.3 Comp. steel AA-4 — — 834.5 36533.2 83.3 698 47.6 29.4 764.4 22.3 86.5 Inv. steel AB-4 1300 5.3 719.627 491.8 32 659.5 32.9 69 784.1 6.1 50.4 Inv. steel AC-4 1049 0.6 898.563 540.1 36.5 717.1 39.5 64.7 782.3 5.9 14.8 Inv. steel AD-4 1264 6.4872.2 100 530.3 104.7 707.5 6.8 47.5 747.7 19.4 12.5 Inv. steel AE-41160 5.0 808.3 55 556.7 67.8 655.3 38.3 72.6 735 28.4 45.8 Inv. steelAF-4 1011 3.4 885.6 60 402.3 24 662.5 45.8 17.2 742.7 36.9 55.7 Comp.steel AG-4 1044 0.4 716.6 98 496.5 55.1 684.3 23.5 67.4 756.7 41.2 22.4Inv. steel AH-4 1286 5.0 696.9 92 444.3 65.3 656.9 44.4 46.8 750.3 28.521.8 Inv. steel AI-4 1054 1.7 754.6 90 452.1 55.4 705.2 51.8 14.6 775.639.6 68   Inv. steel AJ-4 1233 5.9 772.2 22 536.2 63.2 675.9 26.3 3756.9 66 23.7 Comp. steel AK-4 1010 5.6 940.5 47 554.4 122.4 689.8 30.827.1 759.8 27 59   Inv. steel AL-4 1199 6.2 846.4 92 557.1 99.3 706.225.6 64 753.3 20.6 99.2 Inv. steel AM-4 1239 3.6 750.5 38 500.7 150688.9 31.9 44.5 780.5 49.5 34.3 Inv. steel AN-4 1256 5.6 956 74 498149.2 690.8 8 74.8 744.3 38.8 76.6 Comp. steel AO-4 1241 6.7 750.8 56454 118.1 657.9 61.5 47.4 747.4 43.6 20   Comp. steel AP-4 1043 4.8800.3 78 417.4 31.9 707 7.8 0.5 766.1 8.1  2.7 Comp. steel AQ-4 1032 2.2793.4 43 559.4 64.8 718 13 51.2 748.4 17.1 15.9 Inv. steel

Table 9 shows the results of measurement and results of evaluation ofthe carbide size, pearlite area ratio, ferrite grain size, Vickershardness, ratio of the number of carbides at the ferrite grainboundaries to number of carbides in the ferrite grains, ratio of maximumcrack length to plate thickness at the vertical wall parts, and impactresistance characteristic in the prepared samples.

TABLE 9 No. of carbides at grain Ratio boundaries/ of Ferrite PearliteNo. of maximum Carbide grain area Vickers carbides crack Impact sizesize rate hardness inside length resistance (μm) (μm) (%) (HV) grains(%) characteristic Remarks AA-2 0.78 23.8 9.1 195.2 13.70 15.6 NGComparative Steel AB-2 0.78 26.2 1.2 108.9 3.56 1.6 OK Invention SteelAC-2 0.74 13.0 0.0 120.8 1.35 3.2 OK Comparative Steel AD-2 0.75 23.210.3 214.7 2.81 18.8 NG Comparative Steel AE-2 0.17 14.7 0.1 121.4 1.265.6 OK Invention Steel AF-2 0.97 18.0 13.1 217.2 8.98 13.5 NGComparative Steel AG-2 0.56 11.5 0.4 122.1 3.35 2.2 NG Comparative SteelAH-2 0.46 21.7 1.4 105.3 1.20 4.7 OK Invention Steel AI-2 0.49 9.4 0.2132.6 6.21 2.0 OK Invention Steel AJ-2 0.44 20.8 1.0 125.0 1.92 4.7 OKInvention Steel AK-2 0.37 10.7 1.2 127.5 5.48 2.0 OK Invention SteelAL-2 0.52 12.1 0.9 137.5 5.69 2.3 OK Invention Steel AM-2 0.81 16.5 1.4116.4 9.20 1.3 OK Invention Steel AN-2 0.75 20.9 0.9 117.0 1.35 5.1 OKInvention Steel AO-2 0.43 10.0 1.5 144.9 1.34 6.9 OK Invention SteelAP-2 0.48 16.4 14.2 228.1 2.16 19.9 NG Comparative Steel AQ-2 0.65 26.70.9 113.1 2.74 2.0 OK Invention Steel AA-3 0.73 17.0 0.1 106.1 9.74 0.8OK Invention Steel AB-3 0.48 14.3 0.7 124.1 3.28 2.3 OK Invention SteelAC-3 1.02 12.3 1.2 125.9 11.41 1.4 OK Invention Steel AD-3 0.36 7.5 1.3147.6 1.18 7.7 OK Invention Steel AE-3 0.65 21.5 0.8 111.5 5.36 1.4 OKInvention Steel AF-3 0.63 23.3 1.5 110.3 2.88 1.9 OK Invention SteelAG-3 0.57 17.0 0.1 110.2 3.90 1.6 OK Invention Steel AH-3 0.79 14.5 1.5115.4 4.28 13.1 NG Comparative Steel AI-3 0.75 25.1 0.6 105.8 1.93 2.1NG Comparative Steel AJ-3 0.21 14.8 1.5 129.8 1.27 6.2 OK InventionSteel AK-3 0.46 9.0 0.8 133.6 0.66 10.6 NG Comparative Steel AL-3 0.657.4 0.5 153.7 0.14 11.7 NG Comparative Steel AM-3 0.83 21.1 8.2 189.85.20 12.7 NG Comparative Steel AN-3 0.51 25.0 1.3 116.2 3.59 1.9 NGComparative Steel AO-3 0.61 15.6 0.4 131.8 5.45 2.1 OK Invention SteelAP-3 0.57 8.8 1.0 146.8 1.11 6.8 OK Invention Steel AQ-3 2.17 20.1 1.4119.5 23.93 21.2 NG Comparative Steel AA-4 0.66 19.7 1.0 100.3 2.07 1.8OK Invention Steel AB-4 0.68 20.7 1.1 115.2 10.50 1.1 OK Invention SteelAC-4 1.04 24.4 0.7 107.4 4.35 1.4 OK Invention Steel AD-4 0.6 13.3 0.9129.2 7.44 1.8 OK Invention Steel AE-4 0.24 12.2 0.5 126.5 1.52 5.4 OKInvention Steel AF-4 0.36 18.3 0.3 115.8 1.62 2.7 OK Comparative SteelAG-4 0.73 19.3 1.1 107.0 7.13 1.1 OK Invention Steel AH-4 0.52 21.9 1.0105.2 5.55 1.2 OK Invention Steel AI-4 0.78 24.1 0.7 106.3 1.74 2.2 OKInvention Steel AJ-4 0.44 22.4 8.4 199.8 2.88 13.7 NG Comparative SteelAK-4 0.45 16.5 0.3 115.9 1.92 4.2 OK Invention Steel AL-4 0.56 12.6 0.4176.3 1.74 5.7 OK Invention Steel AM-4 0.77 22.3 0.1 109.1 3.38 1.7 OKInvention Steel AN-4 0.34 20.2 1.3 119.5 2.05 2.6 NG Comparative SteelAO-4 2.18 18.0 1.2 132.1 4.03 14.0 NG Comparative Steel AP-4 1.65 24.71.0 123.2 7.68 1.6 OK Comparative Steel AQ-4 0.52 16.4 1.1 122.2 1.684.9 OK Invention Steel

In Comparative Steel AC-2, the finish hot rolling temperature is low andthe productivity is low. In Comparative Steel AN-4, the finish hotrolling temperature is high, scale flaws form at the surface of thesteel plate and cracks form from the flaw parts when impact load wasgiven after cold forging and carburizing, quenching, and tempering, andthe impact resistance characteristic falls.

In Invention Steel AB-3, the cooling rate at the ROT is slow, so a dropin productivity and formation of scale flaws are invited. In InventionSteels AJ-3 and AD-4, the cooling rate at the ROT is 100° C./sec, theoutermost layer part of the steel plate is excessively cooled, and finecracks are formed at the outermost layer part.

In Comparative Steel AN-3, the coiling temperature is low, large amountsof bainite, martensite, and other low temperature transformed structuresare produced resulting in embrittlement, fractures frequently occur atthe time of pay out of the hot rolled coil, and the productivity falls.Furthermore, at the sample taken from the cracked slab, the cold forgingand impact resistance characteristic after carburizing, quenching, andtempering are inferior.

In Comparative Steel AH-3, the coiling temperature is high, bulkypearlite of the lamellar spacing is formed in the hot rolled structure,needle-shaped coarse carbides are high in thermal stability, and evenafter two-stage step type annealing, the above carbides remain in thesteel plate, so the cold forgeability is low.

In Comparative Steel AF-4, the heating rate in the first stage annealingof the two-stage step type annealing is slow, so the productivity islow. In Comparative Steel AG-2, the heating rate in the first stageannealing is fast, so the difference in temperature between the insidepart and outer circumferential part of the coil becomes larger,scratches and seizing due to the difference in heat expansion occur, andthe cold forging and impact resistance characteristic after carburizing,quenching, and tempering fall.

In Comparative Steel AA-2, the holding temperature in the first stageannealing (annealing temperature) is low, the coarsening of the carbidesat the Ac1 point or less is insufficient, the thermal stability of thecarbides becomes insufficient, the carbides remaining at the time of thesecond stage annealing decrease, the pearlite transformation cannot besuppressed in the structure after gradual cooling, and the coldforgeability falls.

In Comparative Steel AM-3, the first stage holding temperature(annealing temperature) is high, austenite is produced during theannealing, the stability of the carbides cannot be raised, and the coldforgeability and impact resistance characteristic after carburizing,quenching, and tempering fall. In Comparative Steel AF-2, the holdingtime in the first stage annealing is short, the stability of thecarbides cannot be raised, and the cold forgeability is low. InComparative Steel AO-4, the holding time in the first stage annealing islong and the productivity is low.

In Comparative Steel AP-4, the heating rate at the second stageannealing in the two-stage step type annealing is slow, so theproductivity is low. In Comparative Steel AI-3, the heating rate at thesecond stage annealing is fast, so the temperature difference betweenthe inside part and the outer circumferential part of the coil becomegreater, scratches and seizing occur due to the large difference in heatexpansion due to transformation, and, when an impact load is given aftercarburizing, quenching, and tempering, fractures occur from the flawparts and the impact resistance characteristics fall.

In Comparative Steel AL-3, the holding temperature in the second stageannealing (annealing temperature) is low, the amount of production ofaustenite is small, it is not possible to increase the number ratio ofcarbides at the ferrite grain boundaries, and the cold forgeabilityfalls. In Comparative Steel AD-2, the holding temperature in the secondstage annealing (annealing temperature) is high, the dissolution ofcarbides during annealing is promoted, and therefore it becomesdifficult to cause the production of carbides at the grain boundariesafter gradual cooling, and the cold forgeability and impact resistancecharacteristics after carburizing, quenching, and tempering fall.

In Comparative Steel AJ-4, the holding time in the second stageannealing is long and dissolution of carbides is promoted, so the coldforgeability is low. In Comparative Steel AQ-3, the cooling rate fromthe second stage annealing to 650° C. is slow so the productivity is lowand coarse carbides are formed in the structure after gradual cooling socracks formed starting from the coarse carbides at the time of coldforging and the cold forgeability dropped. In Comparative Steel AP-2,the cooling rate from the second stage annealing to 650° C. was slow,pearlite transformation occurred at the time of cooling, and thehardness increased, so the cold forgeability fell.

Here, FIG. 3 shows the relationship among the ratio of the number ofcarbides at the grain boundaries to the number of carbides in thegrains, and the crack length and impact resistance characteristics ofcold forging test pieces after carburizing, quenching, and tempering.

From FIG. 3, it will be understood that if the number ratio (=number ofcarbides at the grain boundaries/number of carbides in grains) exceeds1, it is possible to keep down the ratio of the length of cracksintroduced by cold forging and possible to obtain excellent impactresistance after carburizing, quenching, and tempering.

Further, FIG. 4 shows another relationship between the ratio of thenumber of carbides at the grain boundaries to the number of carbides inthe grains and the crack length of cold forging test pieces and impactresistance characteristic after carburizing, quenching, and tempering.FIG. 4 is a view showing that it is possible to keep down crack lengtheven in steel plate to which additional elements are added.

From FIG. 4, it will be understood that even if adding a suitable rangeof elements to steel plate, if the number ratio (=number of carbides atthe grain boundaries/number of carbides in grains) exceeds 1, it ispossible to keep down the ratio of the length of cracks introduced bycold forging and possible to obtain excellent impact resistance aftercarburizing, quenching, and tempering.

INDUSTRIAL APPLICABILITY

As explained above, according to the present invention, it is possibleto provide low carbon steel plate excellent in cold forgeability andimpact resistance characteristic after carburizing, quenching, andtempering and a method of production of the same. The steel plate of thepresent invention is, for example, suitable as a material when forming apart by cold forging such as plate working to obtain a high cycle gearor other part, so the present invention has high industrialapplicability.

REFERENCE SIGNS LIST

-   -   1. disk-shaped test material    -   2. cup-shaped test material    -   3. crack    -   4. sample    -   5. dropping weight    -   L. maximum length of crack

1. A steel plate being a low carbon steel plate having a chemical composition consisting of, by mass %, C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to 1.00%, Al: 0.001 to 0.10%, Cr: 0.50 to 2.00%, Mo: 0.001 to 1.00%, P: 0.020% or less, S: 0.010% or less, N: 0.020% or less, O: 0.020% or less, Ti: 0.010% or less, B: 0.0005% or less, Sn: 0.050% or less, Sb: 0.050% or less, As: 0.050% or less, Nb: 0.10% or less, V: 0.10% or less, Cu: 0.10% or less, W: 0.10% or less, Ta: 0.10% or less, Ni: 0.10% or less, Mg: 0.050% or less, Ca: 0.050% or less, Y: 0.050% or less, Zr: 0.050% or less, La: 0.050% or less, Ce: 0.050% or less, and a balance of Fe and impurities, wherein a metal structure of the steel plate has: a carbide grain size of 0.4 to 2.0 μm; a pearlite area ratio of 6% or less; and a ratio of the number of carbides at ferrite grain boundaries to the number of carbides inside ferrite grains of more than 1, and the steel plate has a Vickers hardness of 100 HV to 180 HV.
 2. A method of production of the steel plate according to claim 1, the method of production comprising: hot-rolling a steel slab with a chemical composition according to claim 1 to obtain a hot rolled steel plate, the hot-rolling in which finish hot-rolling is completed in a 650° C. to 950° C. temperature region; coiling the hot rolled steel plate at 400° C. to 600° C.; pickling the coiled hot rolled steel plate, and subjecting a first stage annealing to the pickled hot rolled steel plate, the first stage annealing in which the pickled hot rolled steel plate is heated, at a heating rate of 30° C./hour to 150° C./hour, up to an annealing temperature of 650° C. to 720° C. and the steel plate is held for 3 hours to 60 hours; then subjecting a second stage annealing to the hot rolled steel plate, the second stage annealing in which the hot rolled steel plate is heated, at a heating rate of 1° C./hour to 80° C./hour, up to an annealing temperature of 725° C. to 790° C. and the steel plate is held for 3 hours to 50 hours as second stage annealing; and cooling the annealed hot rolled steel plate to 650° C. at a cooling rate of 1° C./hour to 100° C./hour. 